Ball milling method to form ceramically coated welding tool

ABSTRACT

A friction stir welding tool comprising a composite of a tungsten-rhenium alloy and hafnium carbide particles, wherein a crystallite size of the tungsten-rhenium alloy is no more than 100 nm, wherein the hafnium carbide particles are dispersed within the tungsten-rhenium alloy, a method of fabricating the friction stir welding tool, and a method of friction stir welding a metal joint using the tool. Various embodiments of the friction stir welding tool, the method of fabricating the tool, and the method of friction stir welding using the tool are provided.

STATEMENT OF FUNDING ACKNOWLEDGEMENT

The funding support provided by King Abdulaziz City for Science andTechnology (KACST) through the Science and Technology Unit at King FahdUniversity of Petroleum and Minerals (KFUPM) under project no. NSTIP11-ADV2130-04 and NSTIP 13-ADV-2155-04 gratefully acknowledged.

BACKGROUND OF THE INVENTION Technical Field

The present invention relates to a friction stir welding tool comprisinga composite of a tungsten-rhenium alloy and hafnium carbide particles,wherein the tungsten-rhenium alloy has a crystallite size of no morethan 100 nm.

Description of the Related Art

The “background” description provided herein is for the purpose ofgenerally presenting the context of the disclosure. Work of thepresently named inventors, to the extent it is described in thisbackground section, as well as aspects of the description which may not20 otherwise qualify as prior art at the time of filing, are neitherexpressly or impliedly admitted as prior art against the presentinvention.

Friction Stir Welding (FSW) process was invented by The WeldingInstitute (TWI) in 1991. See W. Konig and A. Neises: ‘Wear mechanisms ofultrahard, nonmetallic cutting materials’, Wear, 1993, 162, 12-21,incorporated herein by reference in its entirety. The process employs aspinning pin tool that produces frictional heat in the welding of theworkpiece. The pin tool is pressed into contact with a seam to bewelded. A typical FSW system is shown in FIG. 1. See M. Senmweldi and C.Metals, “ADVANCES IN TOOLING MATERIALS FOR FRICTION,” pp. 1-11, 1991,incorporated herein by reference in its entirety. The base metal heatsup due to the rubbing of tool faces by visco-plastic dissipation ofmechanical energy at high strain rates. See Y. Gao, and R. H. Wagoner,“A simplified model for heat generation during the uniaxial tensiletest”, Metallurgical Transactions, 18A: 1001-1009, 1987; H. C. Braga,and R. A. Barbosa, “Simulation of the increase in temperature due toadiabatic heating in hot deformation processes”, Proceedings of 47thBrazilian Association of Metallurgy and Materials (ABM) AnnualConference, pp. 441-457, ABM, 1992; A. C. Nunes, E. L. Jr., Bernstein,and J. C. McClure, “A rotating plug model for friction stir welding”,81st American Welding Society Annual Convention. Chicago, Ill.; Apr.26-28, 2000, each incorporated herein by reference in their entirety.When the heat of the workpiece reaches about 80% of its melting point itbecomes soft and easy to form joining.

FSW has been used for welding low melting point materials such asaluminum. Microstructural examination of aluminum alloys joined by FSWexhibits several zones in the weld or bead. These are Stir Zone (SZ),Thermo-Mechanical Heat Affected Zone (TMAZ) around the nugget and HeatAffected Zone (HAZ). See L. E. Murr, G. Liu, and J. C. McClure, “A TEMstudy of precipitation and related microstructures in friction stirwelded 6061 aluminum” Journal of Materials Science, 33: 1243-1251, 1998;C. G. Rhodes, M. W. Mahoney, W. H. Bingel, R. A. Spurling, and C. C.Bampton, “Effects of friction stir welding on microstructure of 7075aluminum” Scripta Materiala, 36(1): 69-75, 1997; M. W. Mahoney, C. G.Rhodes, J. G. Flintoff, W. H. Bingel, and R. A. Spurling, “Properties offriction stir welded 7075 T651 aluminum” Metallurgical and MaterialsTransactions 29A: 1955-1964, 1998, each incorporated herein by referencein their entirety. Stir zone experiences the highest strain rates andconsequently higher temperatures. See T. Weinberger, N. Enzinger, and H.Cerjak, “Microstructural and mechanical characterisation of frictionstir welded 15-5PH steel” Sci. Technol. Weld. Join, vol. 3, no. 14, pp.210-215, 2009, incorporated herein by reference in its entirety. Thesezones are related to thermomechanical cycle during FSW of aluminum metaland alloys.

Although the FSW process has initially been developed for joiningnon-ferrous materials such as aluminum, by using suitable tool materialsthe use of the process has been extended to harder and higher meltingpoint materials such as steels, titanium alloys and copper. Recently, aconsiderable attention has been given to FSW of high melting temperaturealloys such as steel, due to the process advantages over conventionalwelding methods, which avoids many problems of fusion welding likeporosity, cracking, and solidification. See T. Weinberger, N. Enzinger,and H. Cerjak, “Microstructural and mechanical characterisation offriction stir welded 15-5PH steel” Sci. Technol. Weld. Join, vol. 3, no.14, pp. 210-215, 2009, incorporated herein by reference in its entirety.

Finding a high strength material as the friction stir welding tool iskey to the success of the process. Recent efforts have been dedicated toproduce cost effective and reusable tools for Friction Stir Welding(FSW) steel and hard alloys; however, most of these efforts needimprovement in tool material development aspects. See S. Park, Y. Sato,H. Kokawa, K. Okamoto, S. Hirano and M. Inagaki: ‘Boride formationinduced by pcBN tool wear in friction-stir-welded stainless steels’,Metall. Mater. Trans. A, 2009 40A, (3), 625-636; P. L. Raffo, “Yieldingand fracture in tungsten and tungsten rhenium alloys” J. Less CommonMet., vol. 2, no. 17, pp. 133-149, 1969; R. Rai, H. K. D. H. Bhadeshia,and T. Debroy, “Review: friction stir welding tools,” Sci. Technol.,vol. 16, no. 4, pp. 325-342, 2011, each incorporated herein by referencein their entirety. FSW of these hard alloys put stringent conditions forthe tool material as tool will be exposed to harsh condition during theprocess.

E. Y. Ivanov et al. investigated the synthesis of nanocrystallinetungsten-rhenium alloy by mechanical alloying. See E. Y. Ivanov, C.Suryanarayana, and B. D. Bryskin, “Synthesis of a nanocrystalline W-25wt % Re alloy by mechanical alloying,” vol. 251, pp. 255-261, 1998,incorporated herein by reference in its entirety. They found thatmechanical alloying of a W-25 wt. % Re powder mixture for 14 hr in ahigh energy mill led to the development of nanocrystalline W—Re alloy.Jonathan et al. investigated the effect of temperature and holding timeon the relative density of W-25% Re mixture during spark plasmasintering and it was found that with the increase of temperature andhold time, the relative density decreases as shown in FIG. 8. SeeJonathan A. Webb Indrajit Charit Cory Sparks Darryl P. Butt Megan FraryMark Carroll, “SPS Fabrication of Tungsten-Rhenium Alloys in Support ofNTR” Fuels Development, Nuclear and Emerging Technologies for Space2011, INL/CON-10-20354, incorporated herein by reference in itsentirety. This is attributed to the diffusion of carbon from graphitedies.

M. A Yar et al. synthesized Nano-crystalline W-1% Y₂O₃ powder by amodified solution chemical reaction of ammonium paratungstate (APT) andyttrium nitrate. See M. a. Yar, S. Wahlberg, H. Bergqvist, H. G. Salem,M. Johnsson, and M. Muhammed, “Spark plasma sintering oftungsten-yttrium oxide composites from chemically synthesizednanopowders and microstructural characterization,” J. Nucl. Mater., vol.412, no. 2, pp. 227-232, May 2011, incorporated herein by reference inits entirety. Spark plasma sintering (SPS) was used to consolidate thepowder at 1100 and 1200° C. for various holding times. It was found thatdispersion of yttrium oxide enhanced the sinterability of W powder ascompared to lanthanum oxide. Park et al. synthesized dense, ultrafineWC-10 wt. % Co tool materials by SPS and shaped the tool to perform FSWon steel. See H. Park, H. Youn, J. Ryu, H. Son, H. Bang, and I. Shon,“Fabrication and mechanical properties of WC-10 wt. % Co hard materialsfor a friction stir welding tool application by a spark plasma sinteringprocess,” vol. 13, no. 6, pp. 705-712, 2012, incorporated herein byreference in its entirety. They also investigated the mechanical andmicrostructural investigation of the tool as well as steel. Luo et al.investigated the mechanical properties of W-3.6Re-.26HfC composite at1700-2980K and found that HfC play very important role in strengtheningthe alloy up to 2960 K due its outstanding thermal stability at veryhigh temperature. See Luo, K. S. Shin, and D. L. Jacobson, “Hafniumcarbide strengthening in a tungsten-rhenium matrix at ultrahightemperatures,” Acta Metall. Mater., vol. 40, no. 9, pp. 2225-2232,September 1992, incorporated herein by reference in its entirety.Mingqui et al. studied the growth behavior of HfC dispersed in the W—Rematrix and investigated the effect of its dispersion on the strength ofthe alloy at temperature above 2200K and they found that from 2200 K to2600K there was little growth of HfC with very slow growth rate. See M.Liu and J. Cowley, “Hafnium carbide growth behavior and its relationshipto the dispersion hardening in tungsten at high temperatures,” Mater.Sci. Eng. A, vol. 160, no. 2, pp. 159-167, February 1993, incorporatedherein by reference in its entirety. Rapid growth occurs after 2600K dueto enhanced diffusion along the grain boundaries. John et al. studiedthe high temperature creep behavior of tungsten-4 wt % rhenium-0.32 wt %hafnium carbide at temperatures ranging from 2200 to 2400 K at 40-70MPa. See J. J. Park, “Creep strength of a tungsten-rhenium-hafniumcarbide alloy from 2200 to 2400 K,” Mater. Sci. Eng. A, vol. 265, no.1-2, pp. 174-178, June 1999, incorporated herein by reference in itsentirety. The stress exponent for secondary stage creep was 5.2.Activation energy for this stage was found to be 594 kJ/mol. Rea et al.consolidated the W-1.3 wt % HfC by hot isostatic pressing. See K. E.Rea, V. Viswanathan, a. Kruize, J. T. M. De Hosson, S. O'Dell, T.McKechnie, S. Rajagopalan, R. Vaidyanathan, and S. Seal, “Structure andproperty evaluation of a vacuum plasma sprayed nanostructuredtungsten-hafnium carbide bulk composite,” Mater. Sci. Eng. A, vol. 477,no. 1-2, pp. 350-357, March 2008, each incorporated herein by referencein its entirety. High resolution transmission electron microscopy(HRTEM) of the consolidated sample is indicated the uniform dispersionof nanosize HfC in the tungsten matrix.

Anhua et al. presented the effects of rhenium concentration on thestrength properties of the W—Re—ThO₂ alloys at high temperatures and theyield strength of the alloys decreased with increasing temperature dueto strengthening effect of Re in the system. See A. Luo, K. Shin, and D.Jacobson, “High temperature tensile properties of W—Re—ThO₂ alloys,”Mater. Sci. Eng. A, 1991, incorporated herein by reference in itsentirety. Liu et al. investigated the effect of micro-size (0.2% Zr)alloying and nano-sized (1% Y₂O₃) oxide dispersion in tungsten and thensintered by SPS. Oxygen at W grain boundaries reacts with Zr to formzirconia. See R. Liu, Z. M. Xie, T. Hao, Y. Zhou, X. P. Wang, Q. F.Fang, and C. S. Liu, “Fabricating high performance tungsten alloysthrough zirconium micro-alloying and nano-sized yttria dispersionstrengthening,” vol. 451, pp. 35-39, 2014, incorporated herein byreference in its entirety.

S. K. Rakhunathan et al. studied the high energy and high rateconsolidation of W and W based composites. See S. K. Raghunathan, C.Persad, and D. L. Bourell, “High-energy, High-rate Consolidation ofTungsten and Tungsten-based Composite Powders,” Mater. Sci., vol. 131,pp. 243-253, 1991, incorporated herein by reference in its entirety. Thepure tungsten compact exhibited a ductile failure of the infilteredcopper matrix. Z. Zak Fang et al. reviewed the synthesis, sintering, andmechanical properties of nanocrystalline cemented tungsten carbide. SeeZ. Z. Fang, X. Wang, T. Ryu, K. S. Hwang, and H. Y. Sohn, “Int. Journalof Refractory Metals & Hard Materials Synthesis, sintering, andmechanical properties of nanocrystalline cemented tungsten carbide—Areview,” Int. J. Refract. Met. Hard Mater., vol. 27, no. 2, pp. 288-299,2009, incorporated herein by reference in its entirety. They alsodiscuss the effect of addition of grain growth inhibitor such as VC onthe grain size. They reported that there was almost no grain growth upto 1100° C. and addition of VC inhibits the grain growth at highertemperatures.

David et al. studied the dislocation density as a result of implantationof ions on W—Re ions. See D. E. J. Armstrong and T. B. Britton, “Effectof dislocation density on improved radiation hardening resistance ofnano-structured tungsten-rhenium,” Mater. Sci. Eng. A, vol. 611, pp.388-393, August 2014, incorporated herein by reference in its entirety.Increase in hardness measured by nanoindentation was attributed to theinteraction between irradiation loop and dislocations. Dongju et al.investigated the effect of milling and sintering on the phases of HfC-Wcomposite. See D. Lee, M. A. Umer, H. J. Ryu, and S. H. Hong, “Theeffect of HfC content on mechanical properties HfC-W composites,” Int.J. Refract. Met. Hard Mater., vol. 44, pp. 49-53, May 2014, incorporatedherein by reference in its entirety. This class of composite was sparkplasma sintered at 1800° C. Ozherelyev et al. discussed the X-RayDiffraction studies of Hf—W alloys. See V. V. Ozherelyev, a. I.Bocharov, a. V. Bondarev, and Y. V. Barmin, “X-ray diffraction study ofatomic structure of Hf—W amorphous alloys,” Int. J. Refract. Met. HardMater., vol. 48, pp. 141-144, January 2015, incorporated herein byreference in its entirety. They found that with the increase of tungstencontents, positions of the peaks shift from right to left suggestingsolid solution formation.

Shuaib et al. reported the results of friction stir welding oftube-tubesheet joints made of steel. See F. A. Al-Badour, N. Merah, A.N. Shuaib, and A. Bazoune, “Experimental Investigation of Friction StirSeal Welding of Tube-Tubesheet” Journal of Pressure Vessel Technology,Vol. 137/011402, 2015, incorporated herein by reference in its entirety.Void defects were reported at the root of some welded regions. Largervoids were observed at the joints having holes without chamfers comparedto those with holes with chamfers. Chung et al. conducted a study of FSW(WC tool) for high carbon steel below and above eutectoid temperature.See C. Y. Dong, F. Hidetoshi, N. Kazuhiro, and N. Kiyoshi, “Frictionstir welding of high carbon tool steel (SK85) below Eutectoidtemperature” Transactions of JWRI, vol. 38, no. 1, pp. 37-41, 2009,incorporated herein by reference in its entirety. These authors reportedthe presence of a mixture of pearlite and cementite structure presentbelow A1 (A1 temperature below which there will be no phasetransformation) whereas all other conditions show martensite pluspearlite structure. In this investigation it was found that below A1with 100 mm/min welding speed and 100 rpm rotation speed, themicrostructure is totally pearlite and cementite whereas when conditionswere changed to 200 mm/min and 400 rpm the microstructures at locationsof joints were different at the top 65% was the martensite whichdecreases to 20% at the bottom.

B. W Ahn et al. investigated the FSW of 409L SS by using a siliconnitride tool. See B. W. Ahn, D. H. Choi, D. J. Kim, and S. B. Jung,“Microstructures and properties of friction stir welded 409L stainlesssteel using a Si 3 N 4 tool,” Mater. Sci. Eng. A, vol. 532, pp. 476-479,2012, incorporated herein by reference in its entirety. The base metal(BM) has hot rolled ferrite grain structure and stir zone (SZ) had anequiaxed ferrite grain structure with a diameter of approximately 50 μm.The equiaxed grain size was formed due dynamic recrystallization.Meshram et al. investigated the mechanical behavior of friction stirwelding of stainless steel performed by PCBN pin tool. See M. P.Meshram, B. K. Kodli, and S. R. Dey, “Friction Stir Welding ofAustenitic Stainless Steel by PCBN Tool and its Joint Analyses,”Procedia Mater. Sci., vol. 6, no. Icmpc, pp. 135-139, 2014, incorporatedherein by reference in its entirety. The base metal was found to have608 MPa UTS whereas Friction stir welded sample showed 630 MPa. The weldbehaved almost similar to base metal. Konkol et al. compared theFriction Stir Welding and Submerged Arc Welding of HSLA-65 Steel. SeeKonkol, “Comparison of Friction Stir Weldments and Submerged ArcWeldments in HSLA-65 Steel,” no. July, pp. 187-195, 2007, incorporatedherein by reference in its entirety. The welding of the 3m length ofHSLA-65 with the refractory alloy tool was done successfully. The W—Repin showed almost no wear or change in length at the completion of theweld.

Zafar et al. investigated the effect of friction stir welding parameterson the weld microstructure of mild steel using a W-25% Re pin tool. SeeZ. Iqbal, A. N. Shuaib, F. Al-badour, N. Merah, and A. Bazoune,“Microstructure and hardness of friction stir weld bead on steel plateusing W-25% Re pin tool,” Proc. ASME 2014 12th Bienn. Conf. Eng. Syst.Des. Anal. ESDA Jun. 25-27, 2014, Copenhagen, Denmark, pp. 1-6, 2014,incorporated herein by reference in its entirety. The effect of tooltravel speed on bead surface finish as well as bead width can beobserved from FIG. 9 where the weld bead is superimposed on the loadprofiles of the pin tool. In the first 30 mm length of the bead, thetool was traveling at 15 mm/min zone (A). Toward the end of zone (A),two surface defects or discontinuities, marked by dashed circles,developed, which caused an increase in the magnitude of the axialwelding force.

Buffa et al. investigated a quantitative analysis of the tool life inFSW of 3 mm thick Ti-6V-4Al titanium alloy sheets by using W-25% Re tooland found it successful in welding the alloy. See G. Buffa, L. Fratini,F. Micari, and L. Settineri, “On the Choice of Tool Material in FrictionStir Welding of Titanium Alloys,” vol. 40, 2012, incorporated herein byreference in its entirety. Lienert et al. studied the FSW for joining ofmild steel by characterizing the process of friction stir welds on mildsteel by using refractory tools. See T. J. Lienert, W. L. Stellwag, Jr.,B. B. Grimmett, and R. W. Warke, “Friction Stir Welding Studies on MildSteel” Welding Research, Supplement to the Welding Journal, January, pp.1-9, 2003, incorporated herein by reference in its entirety. Park et al.synthesized dense, ultrafine WC-10 wt. % Co tool materials by SPS andshaped the tool to perform FSW on steel. This tool was used to frictionstir weld the low carbon steel sheet of thickness 2 mm. No visibledefects were present in the nugget of weld. Stir zone is the regionwhere base material comes in direct contact with the tool and heataffected zone is the region where grain growth can be found.

Thompson et al. evaluated the diffusional wear of three differentTungsten base tool with same geometry for the FSW of steel and titaniumalloys. See Thompson, B., Babu, S. S., 2010, Tool degradationcharacterization in the friction stir welding of hard metals, WeldingJournal (Miami, Fla.), 89 (12): 256s-261, incorporated herein byreference in its entirety. They found that Tungsten-Titanium workpiececross diffusion was more rapid than Tungsten-steel cross diffusion.However, microstructural investigation showed that tungsten from thetool was diffused into the steel. They also reported that W—Re andW—Re—HfC tools showed minimal tool degradation. Barnes et al.investigated the effect of tool material on developed microstructure inFS welded HSLA-65 steel. See Barnes, S. J., Bhatti, A. R., Steuwer, A.,Johnson, R., Altenkirch, J., & Withers, P. J. (2012). Friction stirwelding in HSLA-65 steel: part I. Influence of weld speed and toolmaterial on microstructural development. Metallurgical and MaterialsTransactions A, 43(7), 2342-2355, incorporated herein by reference inits entirety. The authors compared the performance of PNCB to W-25% Re,and they found that excessive level of abrasive wear occurred on theW—Re tool as compared with PNCB tool, and it found to be increasing withtool temperature.

FSW tool wear occurs when it passes through the workpiece. The reductionin yield strength of tool may happen due to high load application andelevated temperature generated during the FSW of harder metal and alloyssuch as steel and titanium alloys. Wear of the tool can be due toabrasion, adhesion or diffusion. Due to the involvement of hightemperature, diffusional wear can play a vital role in the tool wear.From thermodynamic point of view, Ellingham diagram can help to find outrelative ability of oxidation at elevated temperature for such metallictools. In most of the cases, tool failures are related to pin ratherthan shoulder as pin has to face more resistance to motion when immersedin the workpiece. Moreover, the pin has lower load bearing capabilitywhen compared to shoulder part which results in higher torsional andbending stresses in the former. Steel and titanium alloys were recentlyfriction stir welded by W—Re alloy tool as discussed earlier.

There are several varieties of tool materials available in the markets,which include: high carbon steels, high speed steels (HSS), and cementedcarbides (WC-based and ceramics, alumina-based, silicon nitride-based,sintered polycrystalline diamond), and sintered polycrystalline cubicboron nitride (PCBN). Diamond and cubic boron nitride (BN) are known assuper hard materials due to their exceptional hardness. So we can saythat going from carbon steels to diamond, the tool material shows anincrease in wear resistance, hardness, plastic deformation resistance,and cost while the thermal shock resistance and ease of fabricationdecrease.

Cubic boron nitride (cubic BN, Knoop hardness 4700 kgfmm⁻²) is notavailable in nature. Synthesis of cubic BN requires the transformationof BN from hexagonal to cubic form at high temperature-high pressure.FIG. 5 compares important mechanical properties of friction and fusionwelds with those of the parent metal. See K. Brookes, “There's more tohard materials than tungsten carbide alone,” Met. Powder Rep., vol. 66,no. 2, pp. 36-37, 39-45, 2011, incorporated herein by reference in itsentirety. Diamond is the hardest materials (Knoop hardness 8000kgfmnm⁻²). The brittleness of some important tool materials such PCBN isa major issue that needs to be addressed.

Commercially pure tungsten (cp-W) is strong at higher temperatures buthas poor toughness at room temperature. Pure tungsten exhibits high wearwhen utilized as a tool material for Friction Stir Welding of steels andtitanium alloys. Exposure of cp-W to temperatures higher than 1473 Kresults in crystallization and brittleness when cooled to roomtemperature. Addition of Re to tungsten lowers the ductile to brittletransition temperature as a result of changing the Peierls stress fordislocation motion. This led to the development of tungsten-rheniumalloys, with W-25 wt-% Re as a candidate material for FSW tools. Steelsand titanium alloys are friction stir welded by W-25 wt-% Re tool. Theweld microstructure can be affected due to interaction with the toolmaterial. Wear of the tool will increase the cost of the tool if thetool has lower yield strength at elevated temperature. The tool materialunder investigation should have high strength, high thermal conductivityand low coefficient of thermal expansion at high temperature. Theinteraction of work piece with the tool is also important at hightemperature. Pin tools made from PCBN and W based alloys have found tobe suitable candidates for FSW of steel and titanium alloys.

PCBN (Polycrystalline Boron Nitride) and other ceramic tool materialssuch as Si₃N₄ are currently being used on commercial scale due to theirhigh hardness and strength at elevated temperatures. However, processingPCBN involves a combination of very high temperatures and pressures. Inaddition, PCBN easily fail during the plunging stage due to its lowtoughness. Tool wear affects not only the tool life but also the weldcharacteristics. FSW of steels with PCBN involves boron and nitrogenpick-up from worn tool leaving the material susceptible to corrosion andpitting. Workpiece may be contaminated with Nitrogen. Nitrogen can alsoreact with oxygen to make detrimental oxides. PCBN has high thermalconductivity (100-250 W m⁻¹ K⁻¹) which results in higher heat loss andlower workpiece temperatures.

Synthesis of the tool materials plays an important role in gauging theperformance of the tool under severe conditions of friction stir weldingof steel. Table 1 provides a comparison of mechanical alloying andsintering parameters for different tungsten base composites mainlyconsolidated by spark plasma techniques whereas Table 2 shows some othertechniques used for the consolidation of tungsten base alloys andcomposites.

TABLE 1 List of tungsten base materials synthesized by varioustechniques Sinter Temp time System (° C.) (min) Technique Ref. Puretungsten 1800  0-15 Plasma K. Cho, L. Kecskes, R. Dowding, B. Schuster,Q. Wei, and R. Z. Valiev, “Nanocrystalline and Ultra-Fine GrainedTungsten for Kinetic Energy Penetrator and Warhead Liner Applications,”Mater. Res., no. June, 2007 W—25Re 2400 180  Cold press E. Y. Ivanov, C.Suryanarayana, and B. D. Bryskin, “Synthesis of a nanocrystalline W— 25wt. % Re alloy by mechanical alloying,” vol. 251, pp. 255-261, 1998 W—Cu— — — Jonathan A. Webb Indrajit Charit Cory Sparks Darryl P. Butt MeganFrary Mark Carroll, “SPS Fabrication of Tungsten-Rhenium Alloys inSupport of NTR” Fuels Development, Nuclear and Emerging Technologies forSpace 2011, INL/CON-10-20354 W—7Ni—0.1Y₂O₃ 1500 30 Cold press S. N.Alam, “Synthesis and characterization of W—Cu nanocomposites developedby mechanical alloying,” Mater. Sci. Eng. A, vol. 433, no. 1-2, pp.161-168, October 2006 W—B₄C 1700 — Cold press F. Jing-lian, L. Tao, C.Hui-chao, and W. Deng-long, “Preparation of fine grain tungsten heavyalloy with high properties by mechanical alloying and yttrium oxideaddition,” J. Mater. Process. Technol., vol. 208, no. 1-3, pp. 463-469,November 2008 W—1%Y₂O₃ 1200 3-5 SPS N. Ünal, “Mechanical means help makebetter tungsten matrix composites,” Met. Powder Rep., vol. 63, no. 10,pp. 28-33, November 2008 W—5Y₂O₃ 1700  3 SPS M. a. Yar, S. Wahlberg, H.Bergqvist, H. G. Salem, M. Johnsson, and M. Muhammed, “Spark plasmasintering of tungsten-yttrium oxide composites from chemicallysynthesized nanopowders and microstructural characterization,” J. Nucl.Mater., vol. 412, no. 2, pp. 227-232, May 2011 W—30 vol % HfC 1850 — SPSY. Kim, K. H. Lee, E.-P. Kim, D.-I. Cheong, and S. H. Hong, “Fabricationof high temperature oxides dispersion strengthened tungsten compositesby spark plasma sintering process,” Int. J. Refract. Met. Hard Mater.,vol. 27, no. 5, pp. 842-846, September 2009 WC—10 wt % Co 1200 12 SPS J.Lee, J.-H. Kim, and S. Kang, “Advanced W—HfC Cermet using In-Situ Powderand Spark Plasma Sintering,” J. Alloys Compd., October 2012

TABLE 2 List of tungsten base composites synthesized by varioustechniques Sinter Temp time System (° C.) (min) Technique Ref.W—3.6Re—0.26HfC — — Arc melted H. Park, H. Youn, J. Ryu, H. Son, H.Bang, and I. Shon, “Fabrication and mechanical properties of WC—10 wt. %Co hard materials for a friction stir welding tool application by aspark plasma sintering process,” vol. 13, no. 6, pp. 705-712, 2012W—3.6Re—0.35HfC — — Arc melted Luo, K. S. Shin, and D. L. Jacobson,“Hafnium carbide strengthening in a tungsten-rhenium matrix at ultrahightemperatures,” Acta Metall. Mater., vol. 40, no. 9, pp. 2225-2232,September 1992 W—4Re—0.32HfC — — Arc melted M. Liu and J. Cowley,“Hafnium carbide growth behavior and its relationship to the dispersionhardening in tungsten at high temperatures,” Mater. Sci. Eng. A, vol.160, no. 2, pp. 159-167, February 1993 W—1.5HfC 1800 240 HIP J. J. Park,“Creep strength of a tungsten-rhenium-hafnium carbide alloy from 2200 to2400 K,” Mater. Sci. Eng. A, vol. 265, no. 1-2, pp. 174-178, June 199990W—7Ni—3Fe 1500 30 MW K. E. Rea, V. Viswanathan, a. Kruize, J. T. M. DeHosson, S. O'Dell, T. McKechnie, S. Rajagopalan, R. Vaidyanathan, and S.Seal, “Structure and property evaluation of a vacuum plasma sprayednanostructured tungsten-hafnium carbide bulk composite,” Mater. Sci.Eng. A, vol. 477, no. 1-2, pp. 350-357, March 2008

Oxidation of the above mentioned materials as a tool is also a seriousconcern during FSW. FIG. 6 shows the high oxidative resistance ofTungsten and Rhenium at high temperature which is the key physicalproperty in the development and the performance of the tool during theservice. The FSW tool failures are mainly attributed to diffusion andwears. Ellingham diagram shown in FIG. 6 provides information about theregions of stability of oxide formation at high temperature. FIGS. 10Aand 10B show oil hardened steel tool used in friction stir welding ofAl6061-20 vol % Al₂O₃ composite. It was noted that wear rate decreases(due to second phase hard particles) after the initial wear.

The challenges of finding a suitable tool material for friction stirwelding of steels and high temperature alloys may be addressed bymanipulating tungsten base alloys and composites. Although pure tungsten(W) has sufficient strength at elevated temperatures, its applicationhas been limited due to very low toughness, particularly at roomtemperature, and also a large amount of wear when used as a toolmaterial for FSW. Since tungsten is also susceptible to embrittlementand recrystallization at temperatures higher than 1200° C., Rhenium (Re)can be added to lower the ductile to brittle transition temperature andincreases the recrystallization temperature. Furthermore, hafniumcarbide particles may increase a microhardness of the tool material.

In view of the forgoing, one objective of the present invention is toprovide a friction stir welding tool made of a composite of atungsten-rhenium alloy and hafnium carbide particles, wherein thehafnium carbide particles are homogenously dispersed within thetungsten-rhenium alloy, and which is fabricated by ball-milling a solidsolution of tungsten-rhenium alloy and hafnium carbide particlesfollowed by spark-plasma-sintering. Spark-plasma-sintering the solidsolution forms a composite, wherein the tungsten-rhenium alloy has acrystallite size in the range of 20 to 100 nm. Another objective of thepresent invention relates to a method of friction stir welding a highstrength metal joint using the friction stir welding tool.

BRIEF SUMMARY OF THE INVENTION

According to a first aspect, the present disclosure relates to afriction stir welding tool, including a composite of a tungsten-rheniumalloy and hafnium carbide particles, wherein the tungsten-rhenium alloyhas crystallites with a crystallite size of no more than 100 nm, andwherein the hafnium carbide particles are dispersed within thetungsten-rhenium alloy in the composite.

In one embodiment, the tungsten-rhenium alloy has a crystallite size ofno more than 80 nm.

In one embodiment, a concentration of rhenium in the tungsten-rheniumalloy is in the range of 20 to 30 wt %, relative to the total weight ofthe tungsten-rhenium alloy.

In one embodiment, a concentration of the hafnium carbide particles isin the range of 1 to 15 vol %, relative to the total volume of thecomposite.

In one embodiment, the friction stir welding tool has a Vickers hardnessof 400 to 550 HV at a temperature of 20 to 30° C.

In one embodiment, the friction stir welding tool has a relative densityof 95% to 99%.

In one embodiment, the friction stir welding tool has a tip and ashoulder, wherein the tip has a cylindrical, a conical, a triangular, ora pyramidal geometry.

In one embodiment, the friction stir welding tool has a tip and ashoulder, wherein the shoulder is cylindrical with a diameter of 5 to 15mm and a height of 1 to 5 mm, wherein the tip is conical with a heightof I to 3 mm, a base diameter of 3 to 10 mm, and a head diameter of 1 to5 mm, and wherein the tip is not threaded.

In one embodiment, the friction stir welding tool has a tip and ashoulder, wherein the tip is threaded.

In one embodiment, the friction stir welding tool further includes atitanium carbide coating, which covers at least a portion of an externalsurface of the friction stir welding tool.

According to a second aspect, the present disclosure relates to a methodof fabricating a composite of a tungsten-rhenium alloy and hafniumcarbide particles, involving i) ball-milling a tungsten-rhenium alloyfor no more than 25 hours to form a first powder, wherein aconcentration of rhenium is in the range of 20 to 30 wt %, relative tothe total weight of the tungsten-rhenium alloy, ii) mixing hafniumcarbide particles with the first powder to form a second powder, whereina concentration of hafnium carbide particles in the second powder is inthe range of 1 to 15 vol %, relative to the total volume of the secondpowder, iii) ball-milling the second powder for no more than 15 hours toform a third powder, iv) spark-plasma-sintering the third powder at atemperature of 1500 to 2000° C. for no more than 10 minutes to form thecomposite of a tungsten-rhenium alloy and hafnium carbide particles,wherein the hafnium carbide particles are dispersed within thetungsten-rhenium alloy, and wherein the tungsten-rhenium alloy has acrystallite size of no more than 100 nm.

In one embodiment, the method further involves extruding the compositeto form a friction stir welding tool that has a tip and a shoulder,wherein the tip has a cylindrical, a conical, a triangular, or apyramidal geometry.

In one embodiment, a crystallite size of the tungsten-rhenium alloy inthe third powder is in the range of 10 to 50 nm.

In one embodiment, the third powder is compacted with a pressure of 40to 60 MPa during the spark-plasma-sintering.

In one embodiment, the tungsten-rhenium alloy is ball-milled in an inertatmosphere with a milling speed of 200 to 300 rpm, wherein aball-to-powder weight ratio is in the range of 6:1 to 10:1.

In one embodiment, the second powder is ball-milled in an inertatmosphere with a milling speed of 100 to 200 rpm, wherein aball-to-powder weight ratio is in the range of 4:1 to 6:1.

According to a third aspect, the present disclosure relates to a methodof friction stir welding a metal joint, involving i) rotating a shaft ofa friction stir welding machine, wherein a friction stir welding tool issecured on the shaft via a tool holder, ii) plunging the friction stirwelding tool into the metal joint to melt at least a portion of themetal joint, iii) moving the shaft along the metal joint to weld themetal joint, wherein the friction stir welding tool includes a compositeof a tungsten-rhenium alloy and hafnium carbide particles, wherein thetungsten-rhenium alloy has a crystallite size of no more than 100 nm.

In one embodiment, the metal joint comprises two steel plates forming abutt joint.

In one embodiment, the shaft is rotated with a rotational speed of 400to 2000 rpm.

In one embodiment, the friction stir welding tool is plunged into themetal joint with a plunging rate of no more than 5 mm/min.

In one embodiment, the shaft is moved along the metal joint with atraverse speed of 15 to 40 mm/min.

In one embodiment, the method of friction stir welding further involvesi) heat-treating the tool holder at a temperature of 800 to 900° C., ii)tempering the tool holder prior to the mounting, wherein the tool holderhas a hardness of 50 to 60 HRC.

The foregoing paragraphs have been provided by way of generalintroduction, and are not intended to limit the scope of the followingclaims. The described embodiments, together with further advantages,will be best understood by reference to the following detaileddescription taken in conjunction with the accompanying drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

A more complete appreciation of the disclosure and many of the attendantadvantages thereof will be readily obtained as the same becomes betterunderstood by reference to the following detailed description whenconsidered in connection with the accompanying drawings, wherein:

FIG. 1 is a schematic of a friction stir welding (FSW) process.

FIG. 2 is a schematic drawing of a planetary ball mill.

FIG. 3 is a schematic of a spark-plasma-sintering process.

FIG. 4 is a schematic of direct current-pulse current flow through theparticles in a spark-plasma-sintering process.

FIG. 5 represents a stress-strain behavior of a FSW, a fusion weld, anda base metal.

FIG. 6 represents the Ellingham diagram for some of metals used as a FSWtool material.

FIG. 7 represents various FSW tool pin geometries.

FIG. 8 represents spark-plasma-sintering of W-25% Re as a function oftime and temperature.

FIG. 9 represents a weld bead quality along a metal joint in FSW at 2000rpm.

FIG. 10A represents wearing of an oil hardened steel tool at 3 mm s⁻¹.

FIG. 10B represents wearing of an oil hardened steel tool at 9 mm s⁻¹.

FIG. 11 represents a particle size analysis of a tungsten-rhenium alloy(W-25% Re) as-received.

FIG. 12 represents a particle size analysis of hafnium carbideparticles.

FIG. 13A is a Field Emission Scanning Electron Microscope (FE-SEM) ofhafnium carbide particles.

FIG. 13B represents an XRD spectrum of the hafnium carbide particles.

FIG. 14A is an FE-SEM micrograph of an alloyed W-25 wt % Re powder.

FIG. 14B is an FE-SEM micrograph of the alloyed W-25 wt % Re powder.

FIG. 14C is an FE-SEM micrograph of the alloyed W-25 wt % Re powder.

FIG. 14D is an FE-SEM micrograph of the alloyed W-25 wt % Re powder.

FIG. 15A is a magnified FE-SEM micrograph of the alloyed W-25 wt % Repowder.

FIG. 15B is a magnified FE-SEM micrograph of the alloyed W-25 wt % Repowder.

FIG. 15C is a magnified FE-SEM micrograph of the alloyed W-25 wt % Repowder.

FIG. 15D is a magnified FE-SEM micrograph of the alloyed W-25 wt % Repowder.

FIG. 16 represents an XRD spectrum of the alloyed W-25 wt % Re powder atdifferent milling times.

FIG. 17 represents a crystallite size of the alloyed W-25 wt % Re powderat different milling times.

FIG. 18 represents a lattice strain of the alloyed W-25 wt % Re powderat different milling times.

FIG. 19A is an FE-SEM micrograph of a milled composite of the alloyedW-25 wt % Re powder and 5vol % hafnium carbide, after 15 hours ofmilling.

FIG. 19B is an FE-SEM micrograph of a milled composite of the alloyedW-25 wt % Re powder and 10 vol % hafnium carbide, after 15 hours ofmilling.

FIG. 20A is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 5vol % hafnium carbide, after 5 hours of milling.

FIG. 20B is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 5vol % hafnium carbide, after 10 hours of milling.

FIG. 20C is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 5vol % hafnium carbide, after 15 hours of milling.

FIG. 21A is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 10vol % hafnium carbide, after 5 hours of milling.

FIG. 21B is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 10 vol % hafnium carbide, after 10 hours of milling.

FIG. 21C is an X-ray mapping of a milled composite of the alloyed W-25wt % Re powder and 10 vol % hafnium carbide, after 15 hours of milling.

FIG. 22 represents XRD spectra of a milled composite of the alloyed W-25wt % Re powder and 5vol % hafnium carbide, after different millingtimes.

FIG. 23A is an optical image of a surface of W-25% Re before etching.

FIG. 23B is an optical image of a surface of W-25% Re after 15 minutesof etching with the Murakami reagent.

FIG. 24A is an optical image of a surface of W-25% Re after 3 minutes ofetching with sodium hydroxide.

FIG. 24B is an optical image of a surface of W-25% Re after 15 minutesof etching with sodium hydroxide that reveals the grains.

FIG. 25 is an XRD spectrum of the sintered composite of the alloyed W-25wt % Re powder and 5vol % hafnium carbide.

FIG. 26A is an optical image of a surface of a sintered W-25% Re after 1second of etching.

FIG. 26B is an optical image of a surface of a sintered W-25% Re after 3second of etching.

FIG. 27A is an FE-SEM micrograph of an etched surface of a sinteredW-25% Re after 3 second of etching.

FIG. 27B is an FE-SEM micrograph of an etched surface of a sinteredW-25% Re after second of etching.

FIG. 27C is an FE-SEM micrograph of an etched surface of a sinteredW-25% Re after 7 second of etching.

FIG. 28A is an FE-SEM micrograph of an un-etched surface of a sinteredsemi-alloyed W-25% Re.

FIG. 28B is an FE-SEM micrograph of an etched surface of a sinteredsemi-alloyed W-25% Re using secondary electron detector.

FIG. 28C is an FE-SEM micrograph of an etched surface of a sinteredsemi-alloyed W-25% Re using back-scatter electron detector.

FIG. 29A is an FE-SEM micrograph of an etched surface of a sinteredcomposite of the alloyed W-25 wt % Re powder and 5vol % hafnium carbideusing secondary electron detector.

FIG. 29B is an FE-SEM micrograph of an etched surface of a sinteredcomposite of the alloyed W-25 wt % Re powder and 5vol % hafnium carbideusing back-scatter electron detector.

FIG. 30A is an FE-SEM micrograph of a surface of a sintered composite ofthe alloyed W-25 wt % Re powder and 5vol % hafnium carbide usingsecondary electron detector.

FIG. 30B is an FE-SEM micrograph of a surface of a sintered composite ofthe alloyed W-25 wt % Re powder and 5vol % hafnium carbide usingback-scatter electron detector.

FIG. 30C is an FE-SEM micrograph of a surface of a sintered composite ofthe alloyed W-25 wt % Re powder and 10 vol % hafnium carbide usingsecondary electron detector.

FIG. 31A is an FE-SEM micrograph of a surface of a sintered composite ofthe alloyed W-25 wt % Re powder and 5vol % hafnium carbide.

FIG. 31B is a mapping diagram of hafnium in a sintered composite of thealloyed W-25 wt % Re powder and 5vol % hafnium carbide.

FIG. 31C is a mapping diagram of carbon in a sintered composite of thealloyed W-25 wt % Re powder and 5vol % hafnium carbide.

FIG. 32 represents XRD spectra of a sintered W-25 wt % Re powder atdifferent sintering temperatures.

FIG. 33 represents a crystallite size of the sintered W-25 wt % Repowder at different sintering temperatures.

FIG. 34 represents a relative density of the sintered W-25 wt % Repowder at different sintering temperatures.

FIG. 35 represents a microhardness of the sintered W-25 wt % Re powderat different sintering temperatures.

FIG. 36 represents a relative density of various sintered composites ata sintering temperature of 1800° C.

FIG. 37 represents a microhardness of various sintered composites at asintering temperature of 1800° C.

FIG. 38 represents a thermal conductivity of the sintered W-25 wt % Repowder at different sintering temperatures.

FIG. 39 represents a thermal conductivity of various sintered compositeshaving different hafnium carbide contents.

FIG. 40 is a schematic of a ball-on-disc setup for a wear test.

FIG. 41A is an FE-SEM micrograph of a worn surface of the sintered W-25wt % Re powder at a sintering temperature of 1500° C.

FIG. 41B is an FE-SEM micrograph of a worn surface of the sintered W-25wt % Re powder at a sintering temperature of 1800° C.

FIG. 41C is an FE-SEM micrograph of a worn surface of the sinteredcomposite of the alloyed W-25 wt % Re powder and hafnium carbide at asintering temperature of 1800° C.

FIG. 42A is an optical illustration of the surface of the sintered W-25wt % Re powder after the wear test, when sintering temperature is 1500°C.

FIG. 42B is an FE-SEM micrograph of the sintered W-25 wt % Re powderthat shows a track of a ball after the wear test, when sinteringtemperature is 1500° C.

FIG. 42C is an FE-SEM micrograph of the sintered W-25 wt % Re powderthat shows surface deformations after the wear test, when sinteringtemperature is 1500° C.

FIG. 42D is an FE-SEM micrograph of the sintered W-25 wt % Re powderthat shows debris formed after the wear test, when sintering temperatureis 1500° C.

FIG. 43A is an optical illustration of the surface of the sintered W-25wt % Re powder after the wear test, when sintering temperature is 1800°C.

FIG. 43B is an FE-SEM micrograph of the sintered W-25 wt % Re powderthat shows a track of a ball after the wear test, when sinteringtemperature is 1800° C.

FIG. 43C is a magnified FE-SEM micrograph of the sintered W-25 wt % Repowder that shows a track of a ball after the wear test, when sinteringtemperature is 1800° C.

FIG. 43D is an FE-SEM micrograph of the sintered W-25 wt % Re powderthat shows surface deformations and debris after the wear test, whensintering temperature is 1800° C.

FIG. 44A is an optical illustration of the surface of the sinteredcomposite of the alloyed W-25 wt % Re powder and hafnium carbide afterthe wear test, when sintering temperature is 1800° C.

FIG. 44B is an FE-SEM micrograph of the surface of the sinteredcomposite of the alloyed W-25 wt % Re powder and hafnium carbide thatshows a track of a ball after the wear test, when sintering temperatureis 1800° C.

FIG. 44C is an FE-SEM micrograph of the surface of the sinteredcomposite of the alloyed W-25 wt % Re powder and hafnium carbide thatshows debris formed after the wear test, when sintering temperature is1800° C.

FIG. 44D is an FE-SEM micrograph of the surface of the sinteredcomposite of the alloyed W-25 wt % Re powder and hafnium carbide thatshows chunk of debris after the wear test, when sintering temperature is1800° C.

FIG. 44E is a magnified FE-SEM micrograph of the debris formed.

FIG. 45 represents wear rates of various composites tested at 5, 10 and15 N for a 10 m sliding distance.

FIG. 46 represents a specific wear rate for various composites.

FIG. 47 represents a coefficient of friction for the sintered W-25 wt %Re powder that is sintered at 1500° C.

FIG. 48 represents a coefficient of friction for the sintered W-25 wt %Re powder that is sintered at 1800° C.

FIG. 49 represents a coefficient of friction for the sintered compositeof the alloyed W-25 wt %/Re powder and 5vol % hafnium carbide that issintered at 1800° C.

FIG. 50A is an image of a friction stir welding tool.

FIG. 50B represents a geometry of the friction stir welding tool.

FIG. 51 is an image of a friction stir welding setup.

FIG. 52A represents a forge force and a spindle torque of the frictionstir welding.

FIG. 52B shows a weld beads, at a shaft rotational speed of 800 rpm, awelding speed of 15 to 40 mm/min, and a plunging depth of 1.65 mm.

FIG. 52C shows a magnified weld beads, at a shaft rotational speed of800 rpm, a welding speed of 15 to 40 mm/min, and a plunging depth of1.65 mm.

FIG. 53A represents a forge force and a spindle torque of the frictionstir welding.

FIG. 53B shows a weld beads, at a shaft rotational speed of 800 rpm, awelding speed of 15 to 40 mm/min, and a plunging depth of 1.65 mm.

FIG. 53C shows a magnified weld beads, at a shaft rotational speed of800 rpm, a welding speed of 15 to 40 mm/min, and a plunging depth of1.65 mm.

FIG. 54A represents a forge force and a spindle torque of the frictionstir welding.

FIG. 54B shows a weld beads, at a shaft rotational speed of 800 rpm, awelding speed of 15 to 40 mm/min, and a plunging depth of 1.65 mm.

FIG. 54C shows a magnified weld beads, at a shaft rotational speed of800 rpm, a welding speed of 15 to 40 mm/min, and a plunging depth of1.65 mm.

FIG. 55A is an FE-SEM micrograph of the surface of the friction stirwelding tool.

FIG. 55B is a magnified FE-SEM micrograph of the surface of the frictionstir welding tool that shows buildup of a base metal on the tool.

FIG. 55C is an image of the friction stir welding tool, after a frictionstir welding process.

FIG. 56 represents a forge force and a spindle torque of the frictionstir welding setup and corresponding weld beads, at a shaft rotationalspeed of 1500 rpm, a welding speed of 15 to 40 mm/min, and a plungingdepth of 1.65 mm.

FIG. 57 represents a forge force and a spindle torque of the frictionstir welding setup and corresponding weld beads, at a shaft rotationalspeed of 2000 rpm, a welding speed of 15 to 40 mm/min, and a plungingdepth of 1.65 mm.

FIG. 58A is an optical image of a weld bead after a friction stirwelding process at a shaft rotational speed of 800 rpm.

FIG. 58B is an optical image of a weld bead after a friction stirwelding process at a shaft rotational speed of 1500 rpm.

FIG. 58C is an optical image of a weld bead after a friction stirwelding process at a shaft rotational speed of 2000 rpm.

FIG. 59A is an optical image of a weld bead after a friction stirwelding process at a shaft rotational speed of 800 rpm.

FIG. 59B is a magnified optical image of a weld bead after a frictionstir welding process at a shaft rotational speed of 800 rpm.

FIG. 59C is a magnified optical image of a weld bead near the stir zone,after a friction stir welding process at a shaft rotational speed of 800rpm.

FIG. 60A is an optical image of a weld bead near the heat affected zone,after a friction stir welding process at a shaft rotational speed of 800rpm.

FIG. 60B is an optical image of a weld bead near the heat affected zone,after a friction stir welding process at a shaft rotational speed of 800rpm.

FIG. 60C is an optical image of a weld bead near the heat affected zone,after a friction stir welding process at a shaft rotational speed of1500 rpm.

FIG. 60D is an optical image of a weld bead near the heat affected zone,after a friction stir welding process at a shaft rotational speed of2000 rpm.

FIG. 61 represents a microhardness across the weld bead at differentshaft rotational speeds.

FIG. 62 represents a microhardness along a depth of the weld bead atdifferent shaft rotational speeds.

FIG. 63 represents an effect of tool travel distance on the tool axialand traverse forces and spindle torque.

FIG. 64 represents an effect of the shaft rotational speed on adiffusion of tungsten from the tool to the weld bead.

FIG. 65A is an image of the friction stir welding tool before a frictionstir welding process.

FIG. 65B is an image of the friction stir welding tool after thefriction stir welding process.

FIG. 66A is an image of a disc-shape friction stir welding tool.

FIG. 66B is an image of a machining setup.

FIG. 66C is an image of a tool holder for a disc-shape friction stirwelding tool.

FIG. 66D is an image of a tool holder for a disc-shape friction stirwelding tool.

FIG. 66E is an image of a cross-section of the tool holder that holdsthe disc-shape friction stir welding tool.

FIG. 67A is an image of the disc-shape friction stir welding tool in thetool holder.

FIG. 67B is an image of a stir spindle.

FIG. 67C shows coolant accessory of the friction stir welding setup.

FIG. 67D is an image of a damaged tool holder.

FIG. 67E is an image of a damaged friction stir welding tool, after thefriction stir welding process.

FIG. 67F is an image of a proper friction stir welding tool, after thefriction stir welding process.

FIG. 68A is an image of a proper friction stir welding tool, before afriction stir welding process.

FIG. 68B is an image of a proper friction stir welding tool, after afriction stir welding process.

FIG. 69A is an image of the disc-shape friction stir welding tool in thetool holder.

FIG. 69B is an image of spot welds on a slipped friction stir weldingtool.

FIG. 69C is an image of a friction stir welding tool that is slippedfrom a tool holder.

FIG. 70A is an optical image of a base metal, after a friction stir spotwelding process at a shaft rotational speed of 800 rpm.

FIG. 70B is an optical image of the base metal near the stir zone, aftera friction stir spot welding process at a shaft rotational speed of 800rpm.

FIG. 70C is a magnified optical image of the base metal near the stirzone, after a friction stir spot welding process at a shaft rotationalspeed of 800 rpm.

FIG. 71A shows the quality of the base metal after a friction stir spotwelding process at a shaft rotational speed of 500 rpm.

FIG. 71B shows the quality of the base metal after a friction stir spotwelding process at a shaft rotational speed of 600 rpm.

FIG. 72 represents a diffusion of tungsten from the tool to the basemetal after a friction stir spot welding process at various shaftrotational speeds.

DETAILED DESCRIPTION OF THE EMBODIMENTS

According to a first aspect, the present disclosure relates to afriction stir welding tool, including a composite of a tungsten-rheniumalloy and hafnium carbide particles, wherein the tungsten-rhenium alloyhas crystallites with an average crystallite size of no more than 100nm, preferably no more than 95 nm, preferably no more than 90 nm,preferably no more than 85 nm. Preferably, an average crystallite sizeof the tungsten-rhenium alloy in the composite is in the range of 20 to100 nm, preferably 30 to 95 nm, preferably 40 to 90 nm, preferably 45 to85 nm, preferably 50 to 80 nm, preferably 60 to 75 nm, preferably 65 to73 nm.

The term “composite” as used herein refers to a solid solution oftungsten-rhenium alloy and hafnium carbide particles after beingspark-plasma-sintered.

The term “crystallite” as used herein refers to a sub-grain structuralelement of the tungsten-rhenium alloy. In one embodiment, the averagecrystallite size of the tungsten-rhenium alloy in the composite iscalculated from peak broadening of corresponding peaks in an XRDspectrum (as shown in FIG. 25) of the composite via the Williamsonmethod, according to the following equation:

${\beta_{tot}\cos \; \theta} = {{C\; ɛ*\sin \; \theta} + \frac{K\; \lambda}{L}}$

wherein β_(tot) is the broadening of a diffraction peak in the XRDspectrum, θ is the angle in which the diffraction peak appeared, Cε isthe crystallite strain, K is the Scherrer constant, λ is the wavelengthof the x-ray, and L is the crystallite size. Accordingly, an averagecrystallite size of the tungsten-rhenium alloy in the composite iscalculated and values are depicted in FIG. 33, with respect tocorresponding sintering temperatures of the composite. See Zak, A.Khorsand, et al. “X-ray analysis of ZnO nanoparticles by Williamson-Halland size-strain plot methods.” Solid State Sciences 13.1 (2011):251-256.

The term “grain” as used herein refers to a structural element of thecomposite that contains tungsten-rhenium alloy. A grain may include aplurality of crystallites. A grain of the alloy in the composite mayhave an average grain size in the range of 0.1 to 10 μm, preferably 0.2to 5 μm, preferably 0.5 to 4 μm, preferably 1 to 3 μm, preferably 1.5 to2 μm. In one embodiment, the average grain size of the composite ismeasured via SEM micrographs from a surface of the composite.Accordingly, Scanning Electron Microscope (SEM) images as presented inFIG. 29 and FIG. 30 provide information about grain size.

In one embodiment, an average crystallite size of the tungsten-rheniumalloy in the composite (i.e. after being spark-plasma-sintered) is fivetimes, preferably four times, preferably three times larger than anaverage crystallite size of the tungsten-rhenium alloy before beingspark-plasma-sintered. Accordingly, the average crystallite size of thetungsten-rhenium alloy may be in the range of 1-60 nm, preferably 2-50nm, preferably 5-40 nm, preferably 10-30 nm before spark-plasmasintering, wherein the average crystallite size of the tungsten-rheniumalloy is in the range of 20 to 100 nm, preferably 30 to 95 nm,preferably 40 to 90 nm, preferably 45 to 85 nm, preferably 50 to 80 nmafter spark-plasma sintering.

In one embodiment, the composite includes substantially equiaxed grainsof the tungsten-rhenium alloy. The term “equiaxed grains” as used hereinrefers to grains that have axes with substantially similar lengths. Inanother embodiment, the composite includes grains of thetungsten-rhenium alloy with an even size distribution. For example, inone embodiment, in any 500 μm⁻, preferably 600 μm³, preferably 700 μm³,preferably 800 μm³, preferably 900 μm³, preferably 1000 μm³ volume ofthe composite, at least 80%, preferably at least 85%, preferably atleast 90%, preferably at least 95%, preferably at least 99% of thegrains have a grain size within 1% standard deviation of the averagegrain size.

In one embodiment, a concentration of rhenium in the tungsten-rheniumalloy is in the range of 20 to 30 wt %, preferably 21 to 29 wt %,preferably 22 to 28 wt %, preferably 23 to 27 wt %, preferably 24 to 26wt %, preferably about 25 wt %, relative to the total weight of thetungsten-rhenium alloy.

In one embodiment, a volumetric concentration of the hafnium carbideparticles in the composite is in the range of 1 to 15vol %, preferably 2to 14vol %, preferably 3 to 13vol %, preferably 4 to 12vol %, preferably5 to 11 vol %, preferably 5 to 10vol %, relative to the total volume ofthe composite. Accordingly, a volume fraction of the hafnium carbideparticles in the composite is in the range of 0.01 to 0.15, preferably0.02 to 0.14, preferably 0.03 to 0.13, preferably 0.04 to 0.12,preferably 0.05 to 0.11, preferably 0.05 to 0.1.

In one embodiment, the hafnium carbide particles are present in thecomposite in a size range of less than 3 μm, preferably less than 2 μm,preferably less than 1 μm, preferably less than 500 nm, preferably lessthan 200 nm, preferably less than 100 nm, preferably less than 50 nm,preferably less than 20 nm, preferably less than 15 nm, preferably lessthan 10 nm. The hafnium carbide particles may have various particulateshape including spherical, cylindrical, disc-shape, star-shape,pyramidal, conical, cubical, etc. In a preferred embodiment, the hafniumcarbide particles are spherical with a diameter of less than 2 μm,preferably less than 1 μm, preferably less than 500 nm, preferably lessthan 200 nm, preferably less than 100 nm, preferably less than 50 nm,preferably less than 20 nm, preferably less than 15 nm, preferably lessthan 10 nm. In another embodiment, the hafnium carbide particles may beagglomerated within the composite, however, the size of agglomerationswhen present is less than preferably less than 2 μm, preferably lessthan 1 μm, preferably less than 500 nm, preferably less than 200 nm,preferably less than 100 nm, preferably less than 50 nm. In oneembodiment, the hafnium carbide particles act as grain growth inhibitorsduring sintering, thus form a composite having an average grain size of0.1 to 5 μm, preferably 0.5 to 4 μm, preferably 1 to 3 μm, preferably1.5 to 2 μm.

In a preferred embodiment, the hafnium carbide particles arehomogenously dispersed within the composite. The term “homogenouslydispersed” as used herein refers to an embodiment where a volumefraction of the hafnium carbide particles in any 500 μm³ or less,preferably 600 μm³ or less, preferably 700 μm³ or less, preferably 800μm³ or less, preferably 900 μm³ or less, preferably 1000 μm³ or lessvolume of the composite falls within 5%, preferably 2%, more preferably1% standard deviation of the mean volume fraction of the hafnium carbideparticles. For example, in one embodiment, the hafnium carbide particlesare homogenously dispersed, wherein in any 500 μm³ or less, preferably600 μm³ or less, preferably 700 μm³ or less, preferably 800 μm³ or less,preferably 900 μm³ or less, preferably 1000 μm³ or less volume of thecomposite, the mean volume fraction of the hafnium carbide particlesfalls within 5%, preferably 2%, more preferably 1% of 0.05.

In one embodiment, the composite has a Vickers hardness of at least 25%,preferably at least 30%, more preferably at least 35% higher than aVickers hardness of the tungsten-rhenium alloy, when hardness ismeasured at a temperature in the range of 20 to 30° C., preferably 22 to28° C., preferably 24 to 26° C., preferably about 25° C. For example, inone embodiment, a Vickers hardness of the tungsten-rhenium alloy is inthe range of 350 to 450 HV, preferably 355 to 430 HV, preferably 360 to420 HV, wherein a Vickers hardness of the composite is in the range of400 to 550 HV, preferably 420 to 540 HV, preferably 440 to 530 HV,preferably 460 to 520 HV, preferably 480 to 500 HV, preferably 490 to498 HV, preferably about 498 HV. Vickers hardness is a measure of aresistance of a solid matter to a permanent deformation when acompressive force is applied.

In another embodiment, the composite includes 3 to 7vol %, preferably 4to 6vol %, preferably about 5vol % of the hafnium carbide particles,wherein the tool has a Vickers hardness of 400 to 500 HV, preferably 420to 480 HV, preferably about 450 HV, at a temperature in the range of 20to 30° C., preferably 22 to 28° C., preferably 24 to 26° C., preferablyabout 25° C. In another embodiment, the composite includes 8 to 12vol %,preferably 9 to 11 vol %, preferably about 10 vol % of the hafniumcarbide particles, wherein the tool has a Vickers hardness of 450 to 550HV, preferably 490 to 500 HV, preferably about 495 HV, at a temperaturein the range of 20 to 30° C., preferably 22 to 28° C., preferably 24 to26° C., preferably about 25° C.

In one embodiment, the composite has a relative density in the range of90% to 99%, preferably 95% to 98%, preferably 96% to 97%. The term“relative density” as used herein refers to a density of the composite,relative to a density of a composite that has a substantially similarcomposition with a porosity of zero. Accordingly, in another embodiment,the composite has a porosity in the range of 0.01 to 1%, preferably 0.05to 0.5%, preferably 0.06 to 0.4%, preferably 0.07 to 0.3%, preferably0.08 to 0.2%, preferably about 0.1%.

Referring now to FIGS. 50A and 50B, in one embodiment, the composite isin a form of a friction stir welding tool that has a tip and a shoulder.In view of that, the tip has a cylindrical, a conical, a triangular, ora pyramidal geometry (as shown in FIG. 7). In a preferred embodiment,the shoulder of the tool is cylindrical with a diameter in the range of5 to 15 mm, preferably 6 to 14 mm, preferably 7 to 13 mm, preferably 8to 12 mm, preferably 9 to 10 mm, preferably about 9 mm; and a height inthe range of 1 to 5 mm, preferably 2 to 4 mm, preferably about 3 mm.Further to this embodiment, the tip of the tool is conical with a heightin the range of I to 3 mm, preferably 1.4 to 2 mm, preferably about 1.6mm; a base diameter in the range of 3 to 10 mm, preferably 4 to 8 mm,preferably 5 to 7 mm, preferably about 6 mm; and a head diameter in therange of 1 to 5 mm, preferably 2 to 4 mm, preferably about 3 mm, whereinthe tip is not threaded. In another embodiment, the friction stirwelding tool has a cylindrical shoulder and a conical tip, wherein thetip is threaded. The composite may be extruded to form the friction stirwelding tool with a tip and a shoulder. The composite may also bewrought to form the friction stir welding tool with a tip and ashoulder, or any other desirable shapes. In a preferred embodiment, thefriction stir welding tool has a disc shape, with a diameter in therange of 5 to 50 mm, preferably 10 to 40 mm, preferably 12 to 30 mm,preferably 15 to 25 mm; and a height in the range of I to 10 mm,preferably 2 to 8 mm, preferably 3 to 6 mm, preferably about 5 mm.

In one embodiment, the composite further includes one or more ceramicparticles selected from a group consisting of aluminum oxide, silica,silicon carbide, aluminum nitride, aluminum titanate, barium ferrite,barium strontium titanium oxide, barium zirconate, boron carbide, boronnitride, zinc oxide, tungsten oxide, cobalt aluminum oxide, siliconnitride, zinc titanate, hydroxyapatite, zirconium oxide, antimony tinoxide, cerium oxide, barium titanate, bismuth cobalt zinc oxide, bismuthoxide, calcium oxide, calcium titanate, calcium zirconate, ceriumzirconium oxide, chromium oxide, cobalt oxide, copper iron oxide, copperoxide, copper zinc iron oxide, dysprosium oxide, erbium oxide, europiumoxide, gadolinium oxide, holmium oxide, indium hydroxide, indium oxide,indium tin oxide, iron nickel oxide, iron oxide, lanthanum oxide,lithium titanate, magnesium aluminate, magnesium hydroxide, magnesiumoxide, manganese oxide, molybdenum oxide, neodymium oxide, nickel cobaltoxide, nickel oxide, nickel zinc iron oxide, samarium oxide, samariumstrontium cobalt oxide, strontium ferrite, strontium titanate, terbiumoxide, tin oxide, titanium carbide, titanium carbonitride, titaniumdioxide, titanium oxide, titanium silicon oxide, ytterbium oxide,yttrium oxide, yttrium aluminum oxide, yttrium iron oxide, and zinc ironoxide. Accordingly, a volume fraction of the ceramic particles presentin the composite is less than 0.02, preferably less than 0.01, morepreferably less than 0.005, relative to the total volume of thecomposite. In one embodiment, the ceramic particles have an averageparticle size of less than 20 nm, preferably less than 15 nm, preferablyless than 10 nm, and a purity of at least 97%, preferably at least 98%,more preferably at least 99%, even more preferably at least 99.5%.

In an alternative embodiment, the composite includes the hafnium carbideparticles and cerium oxide particles, wherein a volume fraction of boththe hafnium carbide particles and the cerium oxide particles are in therange of 1 to 15vol %, preferably 2 to 14vol %, preferably 3 to 13vol %,preferably 4 to 12vol %, preferably 5 to 11 vol %, preferably 5 to 10vol %, relative to the total volume of the composite.

In a preferred embodiment, the friction stir welding tool furtherincludes a titanium carbide coating with a thickness in the range of 10μm to 5 mm, preferably 20 μm to 4 mm, preferably 30 μm to 3 mm,preferably 40 μm to 2 mm, preferably 50 μm to 1 mm, preferably 60 μm to900 μm, preferably 70 μm to 800 μm, preferably 80 μm to 700 μm,preferably 90 μm to 600 μm, preferably 100 lm to 500 μm. The titaniumcarbide coating may cover “at least a portion of an external surface ofthe friction stir welding tool”, which as used herein, refers to thetitanium carbide coating covering at least 50%, preferably at least 60%,preferably at least 70%, preferably at least 80%, preferably at least90%, preferably at least 95%, preferably at least 99% of an externalsurface of the friction stir welding tool. In a preferred embodiment,the friction stir welding tool has a cylindrical shoulder and a conicaltip, wherein a titanium carbide coating covers an entire surface area ofthe tip.

The composite may be wrought, machined, or extruded to form a materialwith different geometries to be utilized in various high-temperatureapplications including, but are not limited to car manufacturing,aerospace, electronics, food, pharmaceutical, medical, sport goods, andthe like.

According to a second aspect, the present disclosure relates to a methodof fabricating a composite of a tungsten-rhenium alloy and hafniumcarbide particles.

The method involves ball-milling a tungsten-rhenium alloy for at least 5hours, preferably at least 10 hours, preferably at least 15 hours,preferably at least 20 hours, but no more than 25 hours to form a firstpowder. Accordingly, a concentration of rhenium in the tungsten-rheniumalloy is in the range of 20 to 30 wt %, preferably 21 to 29 wt %,preferably 22 to 28 wt %, preferably 23 to 27 wt %, preferably 24 to 26wt %, preferably about 25 wt %, relative to the total weight of thetungsten-rhenium alloy. Preferably, the tungsten-rhenium alloy may beball-milled in an inert atmosphere, provided by argon, helium, neon,and/or nitrogen, at room temperature (i.e. a temperature of 20 to 30°C., preferably 22 to 28° C., preferably 24 to 26° C., preferably about25° C.), and atmospheric pressure (i.e. a pressure of about 1 atm).

Alternatively, the tungsten-rhenium alloy is ball-milled in an aqueousmedia, for example, in deionized water. In a preferred embodiment, thetungsten-rhenium alloy is ball-milled in a planetary ball-millingmachine with a tungsten carbide ball and a tungsten carbide vial, androtated with a rotational speed of 200 to 300 rpm, preferably 220 to 280rpm, more preferably about 250 rpm. In some embodiments, said vial andball are not made of steel or any other iron-containing alloys toprevent contamination of the tungsten-rhenium alloy, although said vialand ball may be made of ceramic materials such as titanium carbide,silicon carbide, etc. Preferably, the vial has a volume of 100 to 5000mL, preferably 150 to 2000 mL, preferably 200 to 1000 mL, preferably 220to 500 mL, preferably about 250 mL, whereas the ball is spherical with adiameter in the range of 1 to 100 mm, preferably 5 to 50 mm, preferablyabout 10 mm. In one embodiment, a plurality of balls may be placed inthe vial for ball-milling the tungsten-rhenium alloy. Preferably, aball-to-powder weight ratio is in the range of 6:1 to 10:1, preferably7:1 to 9:1, more preferably about 8:1. In the embodiment where aplurality of balls is used, the ball-to-powder weight ratio iscalculated by dividing the total weight of the balls by the total weightof the powder.

The method further involves mixing hafnium carbide particles with thefirst powder to form a second powder. Accordingly, a volumetricconcentration of hafnium carbide particles in the second powder is inthe range of 1 to 15vol %, preferably 2 to 14vol %, preferably 3 to13vol %, preferably 4 to 12vol %, preferably 5 to 11 vol %, preferably 5to 10vol %, relative to the total volume of the second powder.Preferably, the hafnium carbide particles are mixed with the firstpowder at room temperature (i.e. a temperature of 20 to 30° C.,preferably 22 to 28° C., preferably 24 to 26° C., preferably about 25°C.), and atmospheric pressure (i.e. a pressure of about 1 atm). Theparticles may be mixed in a non-oxidizing environment (e.g. in an inertatmosphere comprising nitrogen, argon, helium, or combinations thereof).In one embodiment, the hafnium carbide particles are mixed with thefirst powder in a roll-milling mixer. The hafnium carbide particles andthe first powder may first be diluted in an aqueous media to form asuspension prior to be mixed with the roll-milling mixer. The aqueousmedia may have a low boiling point, preferably less than 70° C., orpreferably less than 60° C., more preferably less than 40° C., so itcould easily evaporate after roll-milling. On the other hand, theaqueous media does not interact with any of the hafnium carbideparticles or the tungsten-rhenium alloy. Exemplary aqueous media mayinclude, but are not limited to chloroform, acetone, methanol, hexane,diethyl ether, tetrahydrofuran, dichloromethane, or combinationsthereof. In one embodiment, the suspension is sonicated prior to bemixed with the roll-milling mixer.

The method further involves ball-milling the second powder for at least5 hours, preferably at least 10 hours, but no more than 15 hours to forma third powder. Accordingly, an average crystallite size of thecrystallites in the third powder may be in the range of 1-60 nm,preferably 2-50 nm, preferably 5-40 nm, preferably 10-30 nm, preferably12-20 nm, preferably 13-15 nm. Preferably, the second powder isball-milled in an inert atmosphere, provided by argon, helium, neon,and/or nitrogen, at room temperature (i.e. a temperature of 20 to 30°C., preferably 22 to 28° C., preferably 24 to 26° C., preferably about25° C.), and atmospheric pressure (i.e. a pressure of about 1 atm).Alternatively, the second powder may be ball-milled in an aqueous media,for example, in deionized water. In a preferred embodiment, the secondpowder is ball-milled in a planetary ball-milling machine with atungsten carbide ball and a tungsten carbide vial and rotated with arotational speed of 100 to 200 rpm, preferably 120 to 180 rpm, morepreferably about 150 rpm. Said vial and ball may be made of ceramicmaterials such as titanium carbide, silicon carbide, etc. Preferably,the vial has a volume of 100 to 5000 mL, preferably 150 to 2000 mL,preferably 200 to 1000 mL, preferably 220 to 500 mL, preferably about250 mL, whereas the ball is spherical with a diameter in the range of 1to 100 mm, preferably 5 to 50 mm, preferably about 10 mm. In oneembodiment, a plurality of balls may be placed in the vial forball-milling the second powder. Preferably, a ball-to-powder weightratio is in the range of 3:1 to 7:1, preferably 4:1 to 6:1, morepreferably about 5:1. In the embodiment where a plurality of balls isused, the ball-to-powder weight ratio is calculated by dividing thetotal weight of the balls by the total weight of the powder.Ball-milling the second powder may homogenously disperse the hafniumcarbide particles within the tungsten-rhenium alloy (as shown in FIGS.20A, 20B, 20C, 21A, 21B, and 21C).

Each of the above mentioned ball-milling steps may alternatively bereferred to as “mechanical alloying (or MA)” in this disclosure, andtherefore these words may be used interchangeably.

The method further involves spark-plasma-sintering the third powder toform the composite of a tungsten-rhenium alloy and hafnium carbideparticles. Preferably, the third powder is spark-plasma-sinteredimmediately after the ball-milling. In one embodiment, the third powderis inserted in a freeze dryer and dried for at least 12 hrs, preferablyat least 24 hours, until all moisture is removed. The dried powder maybe placed in a desiccator before spark-plasma-sintering. In oneembodiment, the third powder may be sieved prior to thespark-plasma-sintering. Accordingly, an average crystallite size of thetungsten-rhenium alloy in the third powder is in the range of 10-30 nm,preferably 11-20 nm, preferably 12-15 nm, preferably about 13 nm. Inanother embodiment, the third powder may by densified via a coldisostatic pressing (CIP) prior to the spark-plasma-sintering.Accordingly, the third powder is loaded into a latex cylindrical vesseland cold-pressed with a hydraulic pressure in the range of 40,000 to60,000 psi, preferably 45,000 to 55,000 psi, preferably about 50,000psi.

In a preferred embodiment, the third powder is spark-plasma-sintered ata temperature in the range of 1500 to 2000° C., preferably 1600 to 1900°C., preferably 1700 to 1850° C., preferably 1750 to 1825° C., preferablyabout 1800° C., for at least 5 minutes, preferably at least 8 minutes,but no more than 10 minutes. In view of that, a heating rate in therange of 60 to 150° C./min, preferably 80 to 120° C./min, preferably 85to 115° C./min, preferably 90 to 110° C./min, preferably 95 to 105°C./min, preferably about 100° C./min is applied to the third powder. Thehigh heating rate, provided by spark-plasma, causes a rapid sintering ofthe third powder, thus forming a sintered composite, wherein acrystallite size of the tungsten-rhenium alloy is no more than 100 nm,preferably no more than 95 nm, preferably no more than 90 nm, preferablyno more than 85 nm, preferably no more than 80 nm. In a preferredembodiment, the third powder is placed in a hollow cylindrical graphitedie having a diameter in the range of 5 to 50 mm, preferably 10 to 40mm, preferably 12 to 30 mm, preferably 15 to 25 mm; and a height in therange of 1 to 10 mm, preferably 2 to 8 mm, preferably 3 to 6 mm,preferably about 5 mm. Alternatively, the third powder is placed in agraphite mold with various shapes to form the preferred shape of thefriction stir welding tool.

In one embodiment, a process for spark-plasma-sintering the third powderis as follows: i) filling the third powder in a graphite mold, ii)installing said mold in a chamber of a discharge plasma sinteringapparatus, iii) creating a vacuum inside the chamber, iv) applying apressure in the range of 30 to 100 MPa, preferably 40 to 60 MPa,preferably 45 to 55 MPa, preferably about 50 MPa, to the third powderinside the graphite mold while concurrently increasing a temperature ofthe mold and the third powder with a heating rate in the range of 60 to150° C./min, preferably 80 to 120° C./min, preferably 85 to 115° C./min,preferably 90 to 110° C./min, preferably 95 to 105° C./min, preferablyabout 100° C./min, until the temperature reaches a final targettemperature in the range of 1500 to 2000° C., preferably 1600 to 1900°C., preferably 1700 to 1850° C., preferably 1750 to 1825° C., preferablyabout 1800° C., v) isothermally maintaining the third powder at thetarget temperature for at least 5 minutes, preferably at least 8minutes, but no more than 10 minutes, vi) cooling the temperature insideof the chamber while maintaining the pressure applied to the thirdpowder inside the mold.

Although the third powder is preferably spark-plasma-sintered in vacuum,it may be spark-plasma-sintered in a non-oxidizing environment (e.g. inthe presence of nitrogen, argon, helium, neon, or combinations thereof).

In one embodiment, the graphite mold has a shape of a tool with ashoulder and a tip, and thus a molded tool with a shoulder and a tip isformed after the spark-plasma-sintering. In view of that, in someembodiments, the tip has a cylindrical, a conical, a triangular, or apyramidal geometry.

In an alternative embodiment, the temperature of the mold and the thirdpowder is increased to the final target temperature in the range of 1500to 2000° C., preferably 1600 to 1900° C., preferably 1700 to 1850° C.,preferably 1750 to 1825° C., preferably about 1800° C. via a stepwiseheating protocol as follows: i) heating the mold and the third powder toa first target temperature of 550 to 650° C., preferably 575 to 625° C.,preferably about 600° C., at a heating rate of 60 to 150° C./min,preferably 80 to 120° C./min, preferably about 100° C./min, andisothermally maintaining the third powder at the first targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; ii) heating the mold and the third powder to a secondtarget temperature of 900 to 1000° C., preferably 950 to 995° C.,preferably about 990° C., at a heating rate of 30 to 80° C./min,preferably 40 to 70° C./min, preferably about 60° C./min, andisothermally maintaining the third powder at the second targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; iii) heating the mold and the third powder to a thirdtarget temperature of 1000 to 1100° C., preferably 1050 to 1095° C.,preferably about 1090° C., at a heating rate of 10 to 80° C./min,preferably 20 to 60° C./min, preferably 40 to 50° C./min, andisothermally maintaining the third powder at the third targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; iv) heating the mold and the third powder to a fourthtarget temperature of 1100 to 1200° C., preferably 1150 to 1195° C.,preferably about 1190° C., at a heating rate of 10 to 80° C./min,preferably 20 to 60° C./min, preferably 40 to 50° C./min, andisothermally maintaining the third powder at the fourth targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; v) heating the mold and the third powder to a fifthtarget temperature of 1200 to 1300° C., preferably 1250 to 1295° C.,preferably about 1290° C., at a heating rate of 10 to 80° C./min,preferably 20 to 60° C./min, preferably 40 to 50° C./min, andisothermally maintaining the third powder at the fifth targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; vi) heating the mold and the third powder to a sixthtarget temperature of 1300 to 1400° C., preferably 1350 to 1395° C.,preferably about 1390° C., at a heating rate of 10 to 80 CJ/min,preferably 20 to 60° C./min, preferably 40 to 50° C./min, andisothermally maintaining the third powder at the sixth targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes; vii) heating the mold and the third powder to a seventhtarget temperature of 1400 to 2000° C., preferably 1700 to 1900° C.,preferably about 1800° C., at a heating rate of 10 to 80° C./min,preferably 20 to 60° C./min, preferably 40 to 50° C./min, andisothermally maintaining the third powder at the seventh targettemperature for 1 to 10 minutes, preferably 2 to 8 minutes, preferablyabout 5 minutes.

After isothermally maintaining the third powder at the seventh targettemperature, the inside of the chamber is cooled to a room temperature(i.e. a temperature of 20 to 30° C., preferably 22 to 28° C., preferably24 to 26° C., preferably about 25° C.), while the pressure applied tothe third powder inside the mold is maintained. Once the temperature isequilibrated at room temperature, a sintered composite is separated fromthe graphite mold. The sintered composite prepared through theabove-mentioned processes has a structure as illustrated in FIG. 50B.

In the preparation processes, a high current having a low voltage pulsedphase is introduced into a gap between the third powder particles by thecurrent applied through the upper and lower electrodes placed on thegraphite mold, and the sintered composite is molded by thermal diffusionand electro-transport caused by high energy of discharge plasmamomentarily generated by spark discharge, pressure and heat caused byelectric resistance of the mold, and electric energy. Also, the pulsedcurrent activation is a direct heating manner in which the currentdirectly flows to the third powder. Heat is generated in the thirdpowder (or other powder samples) at the same time when the current isapplied to the mold, and a temperature difference between an inside ofthe powder and an outside thereof is relatively small, and also it ispossible to minimize a thermal activation action generated in thesintering process, due to a relative low temperature and a shortsintering time. Preferably, by spark-plasma-sintering a composite oftungsten-rhenium alloy and hafnium carbide particles, it may be possibleto achieve a composite with a relative density of at least 95%,preferably at least 96%, preferably at least 97%, preferably at least98%, preferably at least 99%/, preferably at least 99.5%, and finecrystallites size in the range of 20 to 100 nm, preferably 30 to 95 nm,preferably 40 to 90 nm, preferably 45 to 85 nm, preferably 50 to 80 nm,which are proper for a friction stir welding tool. Furthermore, byspark-plasma-sintering a composite of tungsten-rhenium alloy and hafniumcarbide particles, it may be possible to form a composite having adiameter and thickness which is 10 to 25 times, preferably 15 to 20times larger than a composite sintered via conventional methods (e.g.pressureless sintering, hot pressing, hot isostatic pressing, etc.). Inaddition, a composite which is formed via the spark-plasma-sintering, asdescribed, may have a higher strength, a higher abrasion resistance, ahigher relative density, a lower relative porosity, and a largercrystallite size, when compared to a composite which is sintered viaconventional methods such as pressureless sintering, hot pressing, hotisostatic pressing, etc. Moreover, the method of producing a compositeof tungsten-rhenium alloy and hafnium carbide particles may be simplerand less expensive compare to the conventional methods, yet moreeffective in producing a composite with relatively small crystallites,wherein hafnium carbide particles are homogenously dispersed within.

In view of that, the friction stir welding tool may be fabricated via asingle step sintering process, and may not need additionalpost-sintering processes such as extrusion and/or machining, because atool shape can be formed during the spark-plasma-sintering by using aproper mold. However, in another preferred embodiment, the sinteredcomposite is extruded to form the friction stir welding tool. Thesintered composite may be coated prior to the extrusion process.“Extrusion” as used herein refers to a process through which objectswith a desired cross-section are produced by pushing a material througha die of the desired cross-section. Preferably, the sintered compositemay be extruded at room temperature (i.e. ° C.), although the sinteredcomposite may also be extruded at an elevated temperature in the rangeof 250-400° C., preferably 275-350° C., more preferably about 300° C.

A relative density of the sintered composite may slightly increase by 1%to 3%, preferably about 2%, after the extrusion process. In oneembodiment, the sintered composite is extruded with an extrusion ratioin the range of 10:1 to 20:1, preferably 12:1 to 18:1, more preferably14:1 to 16:1. Extrusion ratio refers to a ratio of a cross-sectionalarea of a material before and after an extrusion. For example, if across-sectional area of a material before an extrusion process is A, anda cross-sectional area of the material after the extrusion processbecomes B, an extrusion ratio of the extrusion process is A:B. Thesintered composite may be coated with a lubricant such as colloidalgraphite, glass powders, silica particles, silicon adhesive, or acombination thereof, before being extruded. The lubricant present on thesurface of the sintered composite may partially or completely be removedafter the sintered composite is extruded. The sintered compositeextrudate may further be wrought to a desired shape. The sinteredcomposite extrudate may also be coated with a ceramic material such astitanium carbide.

According to a third aspect, the present disclosure relates to a methodof friction stir welding a metal joint. The method involves mounting afriction stir welding tool on a tool holder. In one embodiment, thefriction stir welding tool has a shoulder and a tip (as shown in FIGS.50A and 50B) with dimensions as described previously, which is securedin the tool holder such that the tool does not slide relative to thetool holder during friction stir welding. In an alternative embodiment,the friction stir welding tool has a disc shape (as shown in FIGS. 66Aand 68A) with dimensions as described previously, which is secured inthe tool holder such that the tool does not slide relative to the toolholder during friction stir welding. The tool may be secured with screwsand/or other means for fastening in the tool holder.

The tool holder is further secured on a shaft of a friction stir weldingmachine via a rotatable coupling. In view of that, the rotatablecoupling includes a male fitting with a plurality of engagement teethand a threaded sleeve, wherein said teeth are configured to engage withthe shaft of the friction stir welding machine as the threaded sleeve isscrewed on the male fitting.

In one embodiment, the tool holder is heat-treated at a temperature of800 to 900° C., preferably 820 to 880° C., preferably 840 to 860° C.,preferably about 850° C., prior to be secured on the friction stirwelding machine. The tool holder may be quenched in oil. Heat-treatingthe tool holder may prevent damage to the tool holder at an extremecondition of welding a high strength metal joint. Preferably, in anotherembodiment, the tool holder is further tempered prior to be secured onthe friction stir welding machine. Tempering an iron-based alloy refersto a process of heat treating the alloy to increase its toughness. Inview of that, the tool holder is heated to a temperature below a meltingpoint of the alloy and kept isothermal for a certain period of time,followed by cooling down the alloy in dry air. Tempering the tool holdermay reduce a hardness of the tool holder to a value in the range of 50to 60 HRC, preferably 52 to 58 HRC, preferably 53 to 55 HRC, preferablyabout 54 HRC.

The method of friction stir welding further involves rotating the shaftwith a rotational speed of 400 to 2000 rpm, preferably 450 to 1500 rpm,preferably 500 to 1200 rpm, preferably 600 to 1000 rpm, preferably 650to 950 rpm, preferably 700 to 900 rpm, preferably about 800 rpm.

The method of friction stir welding further involves plunging thefriction stir welding tool into the metal joint to melt at least aportion of metals and to weld the metal joint. The term “plunging thefriction stir welding tool into the metal joint” as used herein refersto a process of penetrating the rotating tool into the metal joint tomelt at least a portion of the joint and to weld the metal joint. In oneembodiment, the tool is plunged into the metal joint with a plungingrate of at least 2 mm/min, preferably at least 3 mm/min, but no morethan 5 mm/min. In a preferred embodiment, the shaft is plunged into themetal joint with a tilted angle in the range of no more than 0.5°,preferably no more than 0.4°, preferably no more than 0.3°, preferablyno more than 0.2°, preferably no more than 0.10. In another preferredembodiment, the shaft is plunged into the metal joint to a penetrationdepth of no more than 3 mm, preferably no more than 2 mm, preferably nomore than 1.6 mm. In another preferred embodiment, the shaft is plungedinto the metal joint with a compressive force in the range of 1 to 20kN, preferably 1.5 to 10 kN, preferably 2 to 5 kN, preferably 2.5 to 3kN.

In some embodiments, the tool is plunged at least once. For example, thetool is plunged to a same penetration depth at least 5 times, wherein atleast 10 cm, preferably at least 20 cm, preferably at least 30 cm, ofthe joint metal is welded in each time of plunging.

The method of friction stir welding further involves moving the shaftalong the metal joint to weld the metal joint together. In oneembodiment, the shaft is moved with a traverse speed of 15 to 40 mm/min,preferably 20 to 35 mm/min, preferably 25 to 30 mm/min. The term“traverse speed” as used herein refers to a speed of the shaft in aone-dimensional translational motion. In a preferred embodiment, theshaft is moved along the metal joint with a traverse force in the rangeof 100 to 1000 N, preferably 150 to 800 N, preferably 200 to 700 N,preferably 250 to 600 N, preferably 300 to 550 N, preferably 350 to 500N, preferably about 400 N. The term “traverse force” as used hereinrefers to the amount of force exerted to the shaft to move the shaftalong the metal joint. The phrase “moving the shaft along the metaljoint” refers to a translational motion, wherein a direction of movingthe shaft is parallel to the metal joint, preferably in one dimension.The phrase “moving the shaft along the metal joint” may also refer to atwo dimensional translational motion. For example, in one embodiment,the shaft is moved along the metal joint, wherein a direction of movingthe shaft is parallel to the metal joint, however, the penetration depthbecomes larger as the shaft is moved along the joint (i.e. twodimensional translational motion in x and z directions).

In one embodiment, the metal joint includes two steel plates forming abutt joint. The term “steel” as used herein refers to an iron-carbonalloy, which may also contain Si and/or Cr, and is described as mild-,medium-, or high-carbon steels according to the percentage of carbonpresent in the alloy. Exemplary steels may include, but are not limitedto carbon steel, Damascus steel, stainless steel, austenitic stainlesssteel, ferritic stainless steel, martensitic stainless steel, surgicalstainless steel, tool steel, high strength low alloy (HSLA) steel,advanced high strength steels, ferrous superalloys, and cast iron.

A friction stir welding tool, as described previously, may weld any ofthe above listed steels for a distance of at least 1 meter, preferablyat least 2 meters, preferably at least 3 meters, preferably at least 5meters, preferably at least 8 meters, preferably at least 10 meters,preferably at least 12 meters, preferably at least 15 meters, preferablyat least 18 meters, preferably at least 20 meters, but no more than 25meters.

In a preferred embodiment, the steel plate is one selected from thegroup consist of ASTM A516 Grade 70 carbon steel, AISI 304 austeniticstainless steel, Ferritic Utility grade Stainless steel DIN 1.4003, andASTM A240 Grade UNS 541003. Accordingly, the steel plate may have athickness of 2 to 50 mm, preferably 3 to 40 mm, preferably 5 to 30 mm.

In another preferred embodiment, a non-oxidizing gas (e.g. argon and/orhelium) is flowed to the welding zone to prevent oxidation of the metaljoints.

In an alternative embodiment, the metal joint includes two plates, witheach being a high strength metal or alloy, e.g. a titanium plate or atitanium-alloy plate.

The foregoing paragraphs have been provided by way of generalintroduction, and are not intended to limit the scope of the followingclaims. The described embodiments, together with further advantages,will be best understood by reference to the following detaileddescription taken in conjunction with the accompanying drawings.

The examples below are intended to further illustrate protocols for thefriction stir welding tool, the method of fabricating the tool, and themethod of friction stir welding using the tool, and are not intended tolimit the scope of the claims.

Example 1

The following examples demonstrate the synthesis of W-25% Renanocrystalline alloy reinforced with Hafnium Carbide HfC by usingmechanical alloying and Spark Plasma Sintering (SPS), which was shown tohave a relatively high hardness and high wear resistance, and it cantolerate the harsh conditions of FSW of steel. One reason for selectinga composite of W-25% Re nanocrystalline alloy reinforced with HafniumCarbide for this investigation is to employ the advantage of superiorhigh temperature properties of these elements (all the components havemelting points more than 3000° C.). The addition of Re will decrease theductile to brittle transition temperature of the tool. NanocrystallineW-25% Re solid solution can be synthesized by mechanical alloying (MA)and Spark plasma sintering (SPS). The homogenous dispersion of HfCparticles inside the W-25% Re will enhance the strengthening effect athigh temperature.

Example 2—Materials & Method

A W-25% Re alloy was chosen as the best candidate for the development ofthe tool. One best mode of synthesis of W-25 wt % Re alloy and W-25 wt %Re+Xvol % HfC composites powders via mechanical alloying has beenoverviewed in this example.

Semi-alloyed W-25-wt. % Re and HfC powders supplied by Rhenium Alloys,USA, were used in this investigation. Nanostructured W-25Re alloy andhomogenous W-25Re-HfC composites containing 5 and 10 vol. % of HfCparticles were produced using MA. The experiments were carried out in aplanetary ball mill (Fritsch Pulverisette, PS, Idar-Oberstein, Germany).Since the use of steel vials and balls introduces Fe contamination,tungsten carbide vials (250 mL in volume) and balls (10 mm in diameter)were used to avoid contamination of the powders. See M. S. Boldrick, E.Yang, C. N. J. Wagner J. Non-Cryst. Solids, 150 (1992), pp. 478-482; andE. Y Ivanov, C Suryanarayana, B. D Bryskin, Synthesis of ananocrystalline W-25 wt. % Re alloy by mechanical alloying, MaterialsScience and Engineering: A Volume 251, Issues 1-2, 15 Aug. 1998, Pages255-261, each incorporated herein by reference in their entirety.Milling was performed in argon inert gas to avoid oxidation of thepowders. In the first stage, the as-received and semi-alloyed W-25-wt. %Re powder was milled for 5, 10, 15, and 25h until a singlenanostructured solid solution was obtained. Milling conditions ofball-to-powder weight ratio of 8:1 and speed of 250 rpm were used.

In the second stage, HfC particles (5 and 10 vol. %) were dispersed inthe obtained nanostructured alloy powder using milling conditions of aspeed of 150 RPM, a ball to-powder weight ratio of 5:1, and a millingtime up to 15 h. A summary of the parameters used in the investigationis shown in Table 3.

TABLE 3 Mechanical alloying and milling parameters Time Speed Powdercomposition (hrs) (RPM) BPR Semi-alloy W—5%Re 5, 10, 15 & 25 250 8:1Fully alloyed W—25%Re + 5, 10 & 15 150 5:1 5 vol % HfC Fully alloyedW—25%Re + 5, 10 & 15 150 5:1 10 vol % HfC

Field Emission Scanning Electron Microscope (FE-SEM), Tescan Lyra-3,equipped with Energy Dispersive x-ray Spectroscopy (EDS) was used toanalyze the as received powders, and mechanically alloyed powders. Thedispersion of HfC particles in the synthesized powders was characterizedusing FE-SEM and x-ray mapping using 20 frames. X-ray diffractionexperiments were carried out using a diffractometer (Bruker D8, USA,with a wavelength λ=0.15405 nm) to characterize phases present in thesamples and evaluate the crystallite size and lattice strain of thetungsten phase.

Example 3—Characterization of as-Received Powders

Particles Size Analysis

As-received powders were characterized for particle size analysis byparticle size analyzer. It was found that both the powders were in thesub-micrometer size as shown in FIG. 11 and FIG. 12 respectively. Theresult shows that partially alloyed W-25% Re powder was in the earlystages of milling.

Field Emission-SEM and XRD Pattern of HfC

FE-SEM image of HfC powder is presented in FIG. 13A. The particles havesub-micron size and various shapes ranging from elongated to irregular.FIG. 13B shows XRD pattern of HfC powder, the compound has a cubiccrystal structure of the NaCl type.

Example 4—Characterization of Mechanically Alloyed Powder

FE-SEM Analysis of Synthesized Alloyed Powder

The as-received semi-alloyed W-25 wt % Re powder was synthesized bymechanical alloying of W and Re pure powders; a FE-SEM image showing themorphology of its particles is presented in FIG. 14A. It is known thatin the course of mechanical alloying particles undergo flattening, coldwelding, fracturing, and rewelding. In early stage of milling, plasticdeformation and cold welding dominate and leads to change in particleshape and increase in particle size, respectively. See M. R. Akbarpour,E. Salahi, F. Alikhani Hesari, A. Simchi, H. S. Kim, Microstructure andcompressibility of SiC nanoparticles reinforced Cu nanocomposite powdersprocessed by high energy mechanical milling, Ceramics International 40(2014) 951-960; and C. Suryanarayana, Mechanical alloying and milling,Prog. Mater. Sci. 46 (2001) 1-184, each incorporated herein by referencein their entirety. As can be noticed, in FIG. 14A, some particles haveflattened shape and large particle size while others are more equiaxedand have small particle size. Since the as received W-25 wt % Re powderwas not fully alloyed, it was further milled for different milling timesup to 25 h to achieve complete solubility and obtain a nanostructuredsolid solution.

Mechanical alloying of the powder for 5 hours decreased the particlesize and transformed the shape of particles from flattened to equiaxed,FIG. 14B at low magnification and FIG. 15B at high magnification. Withthe increase in milling time to 15 and 25 hours, the ability of theparticles to strain harden decreases. Therefore, fracturing of particlesbecame significant and leads to a decrease in particle size, FIGS. 14Cand 14D at low magnification and FIGS. 15C and 15D at highmagnification. This is in agreement with published work, where it wasreported that mechanical alloying of W-25% Re powder mixture startingfrom W and Re particles with 10 and 30 μm in size deceased the particlesize and led to the formation of rounded powder particles.

XRD Analysis of Synthesized Alloyed Powder

XRD spectra of W-25 wt % Re powder, mechanically alloyed for differenttimes, are shown in FIG. 16. Analysis of XRD spectrum of the as-receivedW-25 wt % Re powder, confirmed that it has body centered cubic crystalstructure. XRD pattern has reflections both from W (bcc, a=0.3167 nm)and Re (hcp, a=0.2781 nm). A peak characteristic of rhenium is presentalong with W—Re solid solution peaks. This indicates that rhenium is notcompletely dissolved in tungsten. XRD spectra of W-25 wt % Re powdermilled for 5, 10, and hours, presented in FIG. 16, shows that the peakcharacteristic of rhenium is still present but its intensity decreased.However, with the increase of milling time to 25 h, this peakdisappeared and only W—Re solid solution peaks are present. Thisindicates complete solubility of rhenium in tungsten and formation of asingle solid solution.

The peak positions was slightly shifted indicating the solid solutionformation due to inter-diffusions of these two elements. Peaks of Wshifted towards lower 2 theta values indicating decrease of latticeparameters. The atomic radius of W (r_(w)=0.1408 nm) is greater than Re(r_(Re)=0.1375) which results in shifting toward higher 2 theta values.Peak broadening was observed in the later stages due to reduction incrystallite size. The equilibrium maximum solid solubility limit of Rein tungsten was reported to be between 24 to 37% Re depending on thepreparation method. See J. M. Dickinson, L. S. Richardson Trans. ASM, 51(1959), pp. 758-771; E. M. Savitskii, M. A. Tylkina, K. B. Povarova,Rhenium alloys, Izv. Nauka, Moskva (1965); Israel Program for ScientificTranslations, Jerusalem, 1970, pp. 139-334; and R. I. Jaffee, C. T.Sims, J. J. Harwood, The effect of rhenium on the fabricability andductility of molybdenum and tungsten, in: 3rd Plansee SeminarProceedings, Pergamon, N.Y., 1959, pp. 380-411, each incorporated hereinby reference in their entirety. Beyond this solubility, the σ phaseprecipitates and leads to failure of the alloys. In this work, the W-25wt % Re solid solution was obtained by MA and formatin of the σ phasewas not revealed. It was reported that MA could lead to the formation ofstable and metastable phases including solid solutions. See C.Suryanarayana, Metals Mater., 2 (1996), pp. 195-209; and C.Suryanarayana, Bibliography on mechanical alloying and milling,Cambridge International Scientific Publishing, Cambridge, 1995, eachincorporated herein by reference in their entirety. The technique wasused to synthesize W-25% Re single-phase solid solution using steel vialand with steel grinding medium and tungsten carbide vials and balls. SeeF. H. Froes, B. D. Bryskin, C. R. Clark, C. Suryanarayana, E. G.Baburaj, Mechanical alloying of W-25 wt. % Re powder, in: B. D. Bryskin(Ed.), Rhenium and Rhenium Alloys, TMS, Warrendale, P A, 1997, pp.569-583, incorporated herein by reference in its entirety.

Mechanical alloying decreased the intensity of the W—Re solid solutionpeaks. This decrease was accompanied with broadening of the peaks. Thisis due to the fact that mechanical milling of metallic powders isusually associated with a decrease in crystallite size and increase inlattice strain. See Y. Waseda, K. Shinoda, E. Matsubara, X-RayDiffraction Crystallography, Springer-Verlag, Berlin Heidelberg, 2011,incorporated herein by reference in its entirety. The XRD data ofmechanically alloyed W-25% Re powder was used to calculate thecrystallite size and lattice strain as reported elsewhere. See Saheb N,Aliyu I K, Hassan S F, Al-Aqeeli N. Matrix Structure Evolution andNanoreinforcement Distribution in Mechanically Milled and Spark PlasmaSintered Al—SiC Nanocomposites. Materials. 2014; 7(9):6748-6767,incorporated herein by reference in its entirety. FIG. 17 showscrystallite size change of the tungsten phase as function of millingtime. The as-received and partially alloyed powder has a crystallitesize of 188 nm. Milling of the powder for 5 hours decreased thecrystallite size to 68 nm. A further increase in milling time to 10 and15 h led a decrease in crystallite size to 28 and 23 nm, respectively.Equation 1 relates the crystallite size D with the peak broadening 0 andwavelength of x-ray used 1. The value of K is usually taken as 0.9.

$\begin{matrix}{\beta_{hld} = {\frac{K\; \lambda}{D} + {4\; ɛ\; \sin \; \theta}}} & {{Equation}\mspace{14mu} 3.1}\end{matrix}$

A final crystallite size of 13 nm was obtained with the increase inmilling time to 25 h. The decrease in crystallite size is believed totake place in three stages. Formation of large number of dislocationswithin shear bands, in the first stage. Recombination of thesedislocations leads to the formation of small angle grain boundaries, inthe second stage. Finally, the orientation of the formed grains becomerandom, in the third stage. In XRD analysis, the crystallites may referto the size of very small grains (sub-grains) that leads to broadeningof the peak.

It is evident from FIG. 17 that the major reduction in crystallite sizeof the W—Re solid solution took place in the first 10 h. The increase inmilling time beyond 10 h only led to very small change in crystallitesize. This behavior is possible in MA because the smaller grains getsaturated with defects and dislocation pile-ups and, hence, thestructure cannot continue to develop the same way as in large-grainedmetals. See A. S. Khan, B. Farrokh, L. Takacs, Mater. Sci. Eng A, 489(2008) 77, incorporated herein by reference in their entirety.Therefore, once the nanocrystalline structure is fully developed,further decrease in crystallite size become very difficult because ofthe large stress needed to deform the nano-grains. Under theseconditions, creation and motion of dislocations become difficult, andthe structure will continue to deform by grain boundary slidingmechanism.

On the other hand, grain size reduction may be hindered by recovery. SeeS. Hwang, C. Nishimura, P. G. McCormick, Mater. Sci. Eng. A, 318 (2001)22, incorporated herein by reference in its entirety. However,mechanical alloying time was extended to 25 h to have completesolubility of Re in W. Mechanical alloying not only decreased theparticle size and crystallite size but also increased lattice strain inthe tungsten phase as presented in FIG. 18. During mechanical alloying,particles experience heavy plastic deformation, which increasesdislocation density and leads to an increase in lattice strain.

FE-SEM Analysis of the Milled Composite Powders

The W-25% Re alloy powder milled for 25 h was mixed with HfC particles(5 and 10 vol. %) and further milled for different milling times up to15 h. FIGS. 19A and 19B show FESEM images of W-25 wt % Re-5HfC and W-25wt % Re-10HfC composite powders mechanically alloyed for 15h. Since theW-25 wt % Re matrix alloy powder was milled for 25 h to obtain ananostructured single solid solution, as presented in FIG. 14D, furthermilling of the composite powders did not bring about significant changein the powders' particle shape and size. Elemental mapping of Hf and Cin the composite powders mechanically milled for 5 and 10 h revealedthat HfC particles remained relatively agglomerated as shown in FIGS.20A and 20B containing 5 vol. % HfC and FIGS. 21A and 21B containing 10vol. % HfC respectively. However, homogenous composite powders withuniform distribution of HfC particles was obtained with the increase inmilling time to 15 h as it can be seen in typical x-ray mapping of Hfand C, in the composite containing 5 vol. % of HfC and 10 vol. % HfC,presented in FIG. 20C and FIG. 21C, respectively.

XRD Analysis of Synthesized Composite Powders

FIG. 22 shows the XRD patterns of the alloy and composites samplecontaining 5 vol % HfC milled for different length of time. XRD patternof monolithic W-25% Re alloy milled for 25 hrs, the composite containingW-25% Re+5vol % HfC milled for 5 h and 15 h is presented in this figure.As the milling for the synthesis of composites were performed by using aball-to-powder ratio of 5:1 with a speed of 150 rpm, it did not resultin the crystallite size reduction as evident from the broadening of thepeaks. There is almost no broadening of peaks which confirms that theobjective of uniformly distributing of second phase HfC was achievedwithout introducing strains in the composite powder.

Example 5—Synthesis and Development of an Experimental Tool

This example deals with the consolidation of milled powders intocylindrical discs by using SPS. Spark plasma sintered specimens werecharacterized by metallography, field emission scanning electronMicroscopy, X-ray Diffraction technique, microhardness, densitydetermination by Archimedes and thermal conductivity. Later on, wearanalysis was performed on these discs and wear morphology was studied byFESEM and optical profilometer.

Methodology

The prepared powders were consolidated using SPS equipment (FCT system,Germany), model HP D 5. More details on the SPS process were reportedelsewhere. See W. M. Thomas, E. D. Nicholas, J. C. Needham, M. G. Murch,P. Templesmith, and C. J. Dawes, International Patent ApplicationPCT/GB92/02203 and GB Patent Application 9125978.8, 1991, incorporatedherein by reference in its entirety. Disc shaped specimens having 10 mmradius were produced with the help of a graphite die. A thermocouple wasinserted through a drilled hole near the graphite die to record thetemperature during the sintering process. In order to reduce thefriction between the specimen powder and wall of the die, a graphitesheet was placed between them.

Compaction pressure of 50 MPa and heating rate of 100° C./min were usedin all sintering experiments. The nanostructured W-25-wt. % Re alloypowder (milled for 25h) was sintered at temperatures of 1500, 1700, and1800° C. for 10 minutes, to determine the suitable sinteringtemperature; then the as-received W-25-wt. % Re powder and compositescontaining 5 and 10 vol. % HfC were sintered at 1800° C. for 10 minutes.Longer duration and very high temperatures were avoided in order tominimize the chances of diffusion of carbon in the sintered samples. SeeShuaib, A. R., Al-Badour, F., & Merah, N. (2015, July). Friction StirSeal Welding (FSSW) Tube-Tubesheet Joints Made of Steel. In ASME 2015Pressure Vessels and Piping Conference (pp. V06BT06A005-V06BT06A005).American Society of Mechanical Engineers, incorporated herein byreference in its entirety. Jonathan et al. investigated the effect oftemperature and holding time on the relative density of W-25% Re mixtureduring spark plasma sintering and it was found that with the increase oftemperature and hold time, the relative density decreases. This isattributed to the diffusion of carbon from graphite dies.

Example 6—Characterization of Consolidated Specimens

Metallography of Spark Plasma Sintered Specimens

A Hewllet Packard power supply devise, model 6216, was used to etchsintered samples. The etching process was performed in one molarconcentrated solution of NaOH for 3 seconds at a voltage of 5 Volt.Tungsten base alloys and composites are generally difficult to etch byconventional etchants. Metallography feasibility of W-25% Re—HfCsintered samples was investigated by using different etchants forrevealing the grain boundaries. Murakami reagent (10 g KOH or NaOH, 10 gpotassium ferricyanide, 100 mL water), Lactic acid+HNO3 and NaOH wereinitially used in this study. Finally, sintered samples wereelectrolytically etched in 1 M NaOH solution. The consolidated sampleswere characterized microstructurally by optical microscopy and FESEM.

Sintered samples were mounted, ground and polished. W-25% Re fullyalloyed sintered at 1800° C. was etched for 1 to 15 min in Murakamireagent. Etching results in the form of optical images are shown inFIGS. 23A and 23B. The sample did not show any evidence of etching.Lactic acid+HNO3 was subsequently utilized for the same duration butthis etchant was also proved ineffective in revealing the grainboundaries. This sample was etched with NaOH and results of etching areshown in FIGS. 24 and 25. The etchants reacted with the grain boundariesafter 15 min but the etching process was still sluggish.

Electrolytic Etching

After achieving some success in revealing grain boundaries, NaOH waschosen for further investigation. Since the etching was sluggish,electrolytic or electrochemical etching was preferred to accelerate theetching process.

An electrolytic etching setup was used for the etching process.Accordingly, one molar concentrated solution of NaOH was prepared. Thepositive terminal of the low voltage direct current power supply wasconnected to the sample. The negative terminal was connected to a steelplate to make it cathode. Mounted samples were drilled small holes atthe back to get connection and a screw was fitted in that hole. Thescrew touched the sample to make a secure connection. The sample wasplaced in the tank facing another plate which was attached to thecathode with a distance of 6 to 8 centimeters between them. Currentflows from the sample to cathode resulting in the etching of the sample.The voltage was kept constant as 5 volt during the etching process.Specimens were etched for short time interval between 1 to 5 seconds andmicrostructural analysis was conducted after each etching step.

FIGS. 26A and 26B show the electrolytically etched optical images of themonolithic W25% Re alloy sintered at 1800° C. FESEM analysis of thesamples conclude that optimized etching time was 3 seconds. Increasingtime lead to over etching.

FE-SEM Analysis of SPSed Specimens

FIGS. 27A, 27B, and 27C show the FE-SEM images of the W-25% Re sinteredat 1800° C. etched for different times. FIG. 27A shows the FESEM imageof the sample etched that was etched for 3 seconds to reveal the grainboundaries. As the time is increased, over etching happened as shown inFIGS. 27B and 27C. A typical FE-SEM image of the microstructure ofsample sintered at 1800° C. for 10 minutes is presented in FIG. 27A.

FIGS. 28A, 28B, and 28C show the FESEM images of the as-receivedpre-alloyed sample sintered at 1800° C. Secondary Electron (SE) and BackSE images of the etched sample shows that average grain size isapproximately 10 μm. For the composite containing vol % HfC, grainboundaries were revealed and HfC was found embedded in the matrix asshown in FIGS. 29A and 29B. This study was carried out at relatively lowtemperature to achieve nanocrystalline matrix reinforced with HfC. Thesintered nanocrystalline monolithic sample has an average grain size of1-3 um. Due to lower temperature and shorter sintering time, thecomposite and fully alloyed specimens will retain theirnanocrystallinity which will be discussed in detail in later stages.

The microstructure of the composite containing 5 vol. % HfC sintered at1800° C. for 10 minutes is presented in FIGS. 29A and 29B. It revealshomogenous distribution of HfC in the matrix as indicated by arrows. Inorder to distinguish between different phases, both SE and BSE modeswere used as shown in FIGS. 29A and 29B, respectively. It can beconcluded that the homogenous dispersion of HfC in W-25 wt % Re-5HfCobtained by mechanical alloying was maintained in the sintered sample.This homogenous dispersion of the HfC particles, in the W-25 wt %Re-5HfC sintered composite, was confirmed through elemental mapping ofHf and C presented in FIGS. 31A, 31B, and 31C.

In addition, to the advantage of SPS in minimizing grain growth, furtherinhibition of grain growth in the composites is attributed to thepresence of HfC particles. It was reported that HfC particles dispersedat the W grain boundaries inhibited the growth of W grains in HfC-Wcomposites. See Dongju Lee, Malik Adeel Umer, Ho J. Ryu, Soon H. Hong,The effect of HfC content on mechanical properties HfC-W composites,Int. Journal of Refractory Metals and Hard Materials 44 (2014) 49-53,incorporated herein by reference in its entirety. Detailed features forthe 5vol % HfC are presented in FIG. 30A (SE mode) and 30B (BSE mode),which reveal the grain boundaries and homogenous dispersion of 5 vol %HfC in the monolithic nanocrystalline alloy. Composite containing 10 vol% HfC showed resistance to metallography (especially during polishing)due to the presence of relatively large amount of HfC contents andmicrostructural details were also not properly revealed as shown in FIG.30C.

XRD Analysis for Consolidated Monolithic Alloy

FIG. 32 shows XRD spectra of W-25 wt % Re alloy sintered at differenttemperatures for 10 minutes. It can be clearly seen that the intensityof the W—Re peaks increased and their broadening decreased with theincrease in sintering temperature because of the increase in crystallitesize. This is in fact the opposite of what occurred during milling. Itis known that if polycrystalline materials are heated and left at hightemperature, grain growth takes place to reduce the excess energyassociated with grain boundaries. See J. M. Tao, X. K. Zhu, R. O.Scattergood, C. C. Koch, Mater. Design, 50 (2013) 22, incorporatedherein by reference in its entirety. The calculated crystallite size ofsintered W—Re is presented in FIG. 33. It is worth reiterating herethat, before sintering, this crystallite size was 13 nm. Sintering at1500° C. increased the crystallite size to 56 nm. The increase insintering temperature to 1700° C. led to the increase in the crystallitesize to 71 nm. Further increase in sintering temperature to 1800° C.resulted in further increase in crystallite size to 80 nm. FIGS. 32 and33 show that the increase in sintering temperature resulted in theincrease in crystallite size; and the higher the temperature, the morethis growth was because of the enhanced diffusion rate. Isothermal graingrowth dependence on temperature and time is generally described usingthe following simple equation:

G ^(n) −G ₀ ^(n) =Kt  Equation 4.1

where G₀ and G are the grain sizes at initial time t₀ and isothermalholding time t, respectively. K is the material's constant that dependson the temperature:

$\begin{matrix}{K = {K_{0}\exp^{(\frac{- Q}{RT})}}} & {{Equation}\mspace{14mu} 4.2}\end{matrix}$

where Q is the activation energy for grain growth, R is the gas constantand T is temperature. Although sintering led to crystallite size growthin all samples, the average crystallite size of the W—Re solid solutionremained in the nanometer range and did not exceed 80 nm. In SPS, it isclaimed that a local high temperature-state is generated momentarilybecause of spark discharge that takes place in the gap or at the contactpoint between particles. This leads to evaporation or melting on thesurface of particles and formation of necks. In addition to thehigh-localized temperature, the applied pressure and current improveheating rates and reduce sintering time and temperature. Therefore,nanopowders might be consolidated using SPS without excessive graingrowth. Crystallite sizes of the tungsten phase in composites containing5 and 10 vol. % HfC, sintered at 1800° C. for 10 minutes, were 71 and 64nm, respectively, compared to the monolithic sample, sintered under thesame conditions, which had a crystallite size of 80 nm. Overall, thecrystallite size of the matrix phase in the sintered composites remainedin the nanometer range and did not exceed 100 nm.

Density and Microhardness

The bulk density of the consolidated samples was measured according tothe Archimedes principle using Metler Toledo balance densitydetermination KIT model AG285. Digital microhardness tester (Buehler,USA) was used to measure the microhardness of the developed materials.The obtained hardness values were the average of 10 readings. Conditionsof a load of 300 gf and a time of 12 s were used in all measurements.

Relative density and hardness of the fully alloyed and consolidated W-25wt % Re is presented in FIGS. 34 and 35. The sample consolidated at1500° C. for 10 minutes had a relative density of 92%. The increase insintering temperature to 1700° C. increased its relative density to96.2%. A further increase in sintering temperature to 1800° C. increasedits relative density to 97.8%. Overall, the relative density increasedwith the increase in sintering temperature because of the enhanceddiffusion rate. See R. M. German, Sintering Theory and Practice. NewYork, John Wiley, (1996), incorporated herein by reference in itsentirety.

Therefore, the higher the sintering temperature, the higher thediffusion rate and the lower the remaining pores. This can be explainedthrough the dependence of density on sintering temperature as follows.

$\begin{matrix}{\rho = {{s\left( \frac{T}{T_{m}} \right)} + g}} & {{Equation}\mspace{14mu} 4.3}\end{matrix}$

where, ρ is the relative density, s is the temperature sensitivity, T isthe sintering temperature, and T_(m) is the melting temperature. See J.E. Garay, “Current-Activated, Pressure-Assisted Densification ofMaterials,” Annual review of materials research, vol. 40, pp. 445-468,2010. ORRÙ R, LICHERI R, MARIO A LOCCI, CINCOTTI A, CAO G, incorporatedherein by reference in its entirety. On the other hand, the externallyapplied pressure contributes to the rearrangement of particles andbreakdown of agglomerates. This leads to the increase in driving forcefor sintering. In addition, in SPS process, spark plasma, spark impactpressure, Joule heating, and an electrical field diffusion effect couldbe generated by the DC pulse discharge. See Consolidation/synthesis ofmaterials by electric current activated/assisted sintering. Mater SciEng R, 2009, 63: 127-287; VISWANATHAN V, LAHA T, BALANI K, AGARWAL A,SEAL S. Challenges and advances in nanocomposite processing techniques.Mater Sci Eng R, 2006, 54: 121-285; ADACHI J, KUROSAKI K, UNO M,YAMANAKA S. Porosity influence on the mechanical properties ofpolycrystalline zirconium nitride ceramics. J Nucl Mater, 2006, 358:106-110; and JIN X, GAO L, SUN J. Preparation of nanostructuredCrl-xTixN ceramics by spark plasma sintering and their properties [J].Acta Mater, 2006, 54: 4035-4041, each incorporated herein by referencein their entirety. The formation of plasma enhances sintering, however,the role of current is still not clear. See KIM Y H, SEKINO T, KUSUNOSET, NAKAYAMA T, NIIHARA, K, KAWAOKA H. Electrical and mechanicalproperties of K, Ca ionic-conductive silicon nitride ceramics [J]. CeramTrans, 2005, 165: 31-38, incorporated herein by reference in itsentirety. It is believed that a local high temperature state momentarilyoccurs in the gap between particles of the powder because of the sparkdischarge. This induces vaporization and melting of the surfaces of thepowder particles, which significantly increases diffusion rate and leadsto higher densification.

The sample consolidated at 1500° C. for 10 minutes had a microhardnessof 360. The increase in sintering temperature to 1700° C. increased itsmicrohardness to 395. A further increase in sintering temperature to1800° C. increased its microhardness to 422. The hardness of a sinteredmaterial mainly depends on its grain size and porosity. Generally, thegrain size d dependence of the yield stress σ_(ys) is described by ageneral expression (Hall-Petch relationship)

σ_(ys)=σ₀ +kd ^(−1/2)  Equation 4.4

where σ₀ is the lattice friction stress, k is a Hall-Petch slope.Vickers hardness of a polycrystalline material can be related to itsyield strength through a simple relationship H_(v)/σ_(ys)≈3. Therefore,the hardness H_(v) can be related to the grain size by

H _(v) =H ₀ +kd ^(−1/2)  Equation 4.5

where H₀ and k are constants. It is clear from the above relationshipthat the increase in the grain size reduces the hardness of a material.However, hardness of the alloy increased despite the increase in thegrain size because of the fact that during sintering pores areeliminated and the density of the material increases. Therefore, thehardness of the material strongly depends on its relative density andthe effect of grain growth will be small. This is more meaningful,specifically with a process such as spark plasma sintering where theheating rate is high, the sintering temperature is low, and thesintering time is short compared to other conventional sinteringprocesses. See K. Rajeswari, U.S. Hareesh, R. Subasri, D. Chakravarty,R. Johnson, Science of Sintering, 42 (2010) 259; V. Pouchly, K. Maca, Y.Xiong, J. Z. Shen, Science of Sintering, 44 (2012) 169; and D. Veljovi,G. VukoviC, I. Steins, E. Palcevskis, P. S. Uskoković, R. Petrović, D.Janaćković, Science of Sintering, 45 (2013) 233, each incorporatedherein by reference in their entirety. This leads only to very marginalgrain growth as explained above, and therefore the hardness is mainlyinfluenced by density. Since higher density and hardness were obtainedat a sintering temperature of 1800° C., all other samples were sinteredat this temperature for 10 min. FIGS. 36 and 37 show the relativedensity and hardness of partially alloyed W-25 wt % Re, fully alloyedW-25 wt % Re, W-25 wt % Re-5HfC composite, and W-25 wt % Re-10HfCcomposites sintered at 1800° C. for 10 min. The partially and fullyalloyed samples had relative densities of 98.2 and 97.8%, respectively.The composites containing 5 and 10 vol. % of HfC processed relativedensities of 96.9 and 96.2%, respectively. The composites containing 5and 10 vol. % of HfC displayed slightly reduced relative density by−P0.92 and −P1.63%, respectively, with respect to the fully alloyedmonolithic alloy; and by −P1.32 and −P2.03%, respectively, with respectto the partially alloyed monolithic alloy.

Compaction of nanostructured powders or nanocomposites reinforced withhard particles is believed to be more difficult than compaction of theirconventional counterparts because of the larger stresses required andthe higher spring back. Therefore, nanostructured green compacts usuallycontain remaining pores, and may not sinter to full density easily. SeeGrácio, J.; Picu, C. R.; Vincze, G.; Mathew, N.; Schubert, T.; Lopes,A.; Buchheim, C. Mechanical Behavior of Al—SiC Nanocomposites Producedby Ball Milling and Spark Plasma Sintering. Metallurgical and MaterialsTransactions A 2013, 44, 5259-5269; Kamrani, S., Razavi Hesabi, Z.,Riedel, R., Seyed Reihani, S. M. Synthesis and characterization ofAl—SiC nanocomposites produced by mechanical milling and sintering,Advanced Composite Materials 2011, 20, 13-27; Candido, G. M.; Guido, V.;Silva, G.; Cardoso, K. R. Effect of the Reinforcement Volume Fraction onMechanical Alloying of AA2124-SiC Composite, Materials Science Forum2012, 660-661, 317-324, each incorporated herein by reference in theirentirety. On the other hand, the addition of a reinforcement usuallylowers the densification of the composite specifically at large volumefraction. As can be clearly seen in FIGS. 36 and 37, the evolution ofHardness of sintered samples, as presented in FIGS. 34 and 35, followedthe same trend as in densification. The partially alloyed sample had amicrohardness of 360. The fully alloyed sample had a microhardness of423. Since the two alloys possessed very close relative densities i.e.98.2 and 97.8%, respectively, the higher hardness of fully alloyedsample is attributed the effect of mechanical alloying because the fullyalloyed sample was milled for an additional 25 h.

The addition of 5% HfC increased the microhardness to 450. Furtherincrease in HfC content to 10 vol. % increased the microhardness to 495.The composites containing 5 and 10 vol. % of HfC possessed improvedhardness by −P11 and −P22%, respectively, with respect to the fullyalloyed monolithic alloy; and by −P25 and −P37.5%, respectively, withrespect to the partially alloyed monolithic alloy. The compositecontaining 10 vol. % of HfC possessed the highest Vickers hardness valueof 495. As for the composites, the increase in hardness can beattributed to the same factors which lead to the increase in thestrength of particle reinforced metal matrix composites. This includesmall grain size of the matrix (Hall Petch theory), presence ofparticles (Orowan strengthening), increase in dislocations' density,load transfer from the matrix to the reinforcement, and strain gradient.See Munoz-Morris, M.; Garcia Oca, C.; Morris, D. An analysis ofstrengthening mechanisms in a mechanically alloyed, oxide dispersionstrengthened iron aluminide intermetallic. Acta. Mater. 2002, 50,2825-2836; Dai, L.; Ling, Z.; Bai, Y. A strain gradient-strengtheninglaw for particle reinforced metal matrix composites. Scripta Mater.1999, 41, 245-252; Chawla, N.; Shen, Y.-L. Mechanical behavior ofparticle reinforced metal matrix composites. Adv. Eng. Mater. 2001, 3,357-370; and Casati, R.; Vedani, M. Metal matrix composites reinforcedby nano-particles-a review. Metals 2014, 4, 65-83, each incorporatedherein by reference in their entirety.

The decrease in grain size leads to the increase in grain boundaries,which has a significant influence on strength. This is because grainboundaries restrict dislocation motion due to the different orientationof adjacent grains and the discontinuity at the highly disordered grainboundary region. In Orowan strengthening, the moving dislocations bowout between particles and yielding takes place when the bowed-outdislocations become semi-circular in shape, after that dislocationsleaves Orowan loops around the particles. These loops hindersdislocation movement, which leads to work hardening. On the other hand,during cooling from sintering temperature to room temperature,geometrically necessary dislocations are formed due to the difference incoefficient of thermal expansion and modulus of elasticity between thereinforcement and the matrix. This leads to strain hardening of thematerial.

Thermal Conductivity

Thermal conductivity analysis for the consolidated specimens werecarried out by Thermal Conductivity Analyzer equipment. Three drops ofwater were used to make a contact between the specimen and the thermalconductivity probe. Thermal conductivity behavior of monolithic W-25 wt% Re alloy sintered at 1500° C., 1700° C. and 1800° C. is presented inFIG. 38. It was found that thermal conductivity was increased with therise of consolidating temperature. The increase in TC (thermalconductivity) can be attributed to the high relative density achieved athigh temperature as discussed in previous section. At high temperature,degree of porosity will be reduced and particle to particle contact willbe more compact and as a result thermal conductivity will be increased.Thermal conductivity of the composites consolidated samples were foundto decrease with the increase of HfC contents as shown in FIG. 39. HfChas a lower thermal conductivity than the monolithic metallic matrix asa result thermal conductivity decreases with the increase of HfCcontents.

Sliding Wear Behavior

Wear characteristics of monolithic W-25 wt % Re alloy sintered at 1500°C. and 1800° C. and W-25% Re-5HfC composite sintered at 1800° C. havebeen investigated in dry sliding conditions against a steel counter faceusing a pin-on-disk equipment as shown in FIG. 40. W-25 wt % Re-10HfCcomposite was also tested under same conditions but it got heated andresulted in extensive wear of the pin so wear behavior of this compositeis not included in this section. Cylindrical discs having diameter of 20mm and a height of 4 mm were prepared from sintered monolithic alloy andthe composite. Specimens were cut into discs for the wear tests andsurface was prepared for metallography analysis. After metallography,specimens were electrolytically etched in 1 M NaOH for 3 seconds toreveal the microstructure. Before the start of each wear experiment, thespecimens and the balls were both rinsed ultrasonically in acetone.Experiments were carried out by using Oscillating Tribometer TestingMachine. Tests were performed according to ASTMG 99 standards with aload of 5-15N. The linear sliding speed was 50 mm/s and the totalsliding distance was 100m with an angular velocity of 120 rpm. The testswere performed at ambient temperature. Wear analysis of worn outsurfaces and wear debris of nine specimens was carried out by usingFE-SEM and optical profilometer. Wear rate and Coefficient of FrictionCOF were calculated from the test data.

Morphology of Worn Surfaces

FIGS. 41A, 41B, and 41C show FE-SEM images of wear tracks for monolithicand composite specimens. It can be observed that wear track profiles aredifferent in size for the alloy SPSed at two different temperatures asshown in FIGS. 41A and 41B, and composite sample as shown in FIG. 41C.Significant improvements in wear resistance, as indicated by thereduction in wear track width, was attained by increasing sinteringtemperature from 1500° C. to 1800° C. This can attributed to a reducedmicrocracking in the materials due improved density, hardness andreduced porosity as explained earlier and reported by other researcher.See B. K. Prasad, K. Venkateswarlu, O. P. Modi, A. H. Yegneswaran, J.Mater. Sci. Lett. 15 (1996) 1773, incorporated herein by reference inits entirety.

Sintering is a thermally activated process controlled mainly bydiffusion. Monolithic sample sintered at 1500° C. shows greater amountof wear which can be attributed to extensive ploughing actions for thepoorly bonded sintered sample. As porosity increases, the wear becomesmore prominent because pores acts as a source of crack nucleation andpropagation leading to excessive sub-surface fracturing.

Wear resistance was further enhanced by the addition of HfC in thematrix as shown in FIG. 41C. The improvement in the wear resistance ofthe composite was attributed to improved microhardness due to theincrease in the volume fraction of the relatively hard second phasereinforcement. See M. Y. Nazmy, in: N. S. Stoloff, V. K. Sikka (Eds.),Physical Metallurgy and Processing of Intermetallic Compounds, Chapman &Hall, New York, 1996, p. 120, incorporated herein by reference in itsentirety.

FIGS. 42A, 42B, 42C, and 42D show the wear morphology of the monolithicalloy sintered at 1500° C. It has gone for extensive wear with a widthof wear track of 820 μm and a depth of 10 μm as indicated by opticalprofile presented in FIG. 42A. The worn out material tends to buildup onthe inner and outer periphery of wear track as depicted in FIG. 42B. Aprominent surface damage with deformed appearance can be observed on thewear track.

The shape of the debris looks like flakes as noticed in FIG. 42C. Tounderstand the wear mechanisms, surface grooves and marks were analyzed.It can be inferred from the micrographs that wear is predominantly anabrasive in nature. The specimen has lower density and has greateramount of porosity due to lower sintering temperature as discussedearlier. Moreover, it has been reported by some researchers that wearresistance of a material, synthesized by powder metallurgy, is verysensitive to the amount of pores present in the sample. See H.-L. Lee,W.-H. Lu, S. L.-I. Chan, Abrasive wear of powder metallurgy Al alloy6061-SiC composites, Wear 159 (1992) 223-231, incorporated herein byreference in its entirety.

FIGS. 43A, 43B, 43C, and 43D show the wear track images and its detailedmorphology for the monolithic sample sintered at 1800° C. It is evidentthat material has been removed due to ploughing action of the ball onthe material. In some areas, the presence of grooves on the worn surfacewas indicative of abrasive wear due to presence of abrasive marks,parallel grooves and stripes on the track. Pockmarks representingregions from which large chunk of material have been removed can also beobserved. FIG. 43A shows optical profile image which illustratesconsiderable reduction of track width from 821 μm to 185 μm whensintering temperature was increased from 1500° C. to 1800° C. Trackdepth was found to be 1.5 μm. The reduction in wear damage at 1800° C.,can be explained on the bases of enhanced diffusion at high temperatureand improved density due to more elimination of porosity at hightemperature.

FIGS. 44A, 44B, 44C, 44D, and 44E show the FESEM images of the wearprofile and morphology of debris of composite specimen. FIGS. 44B and44C show that large areas were covered with wear debris. These smallsized wear debris have been removed from the material by adhesive wearas shown in FIGS. 44D and 44E. The addition of HfC in the monolithicalloy has reduced the width of the track from 185 μm to 145 μm and depthfrom 1.5 μm to of 1.1 μm as shown in FIG. 44A. The higher wearresistance of the composite can be attributed to its higher values ofhardness as compared to the monolithic alloy. To further understand thedifference in wear behavior between alloy and composite, debris andsliding configuration were observed. There are no marks of abrasive wearon the wear track as depicted for alloy sample.

Each of the debris has a shape of a rounded chip. The wear mechanism canbe identified as adhesion. Continuous sliding friction between HfC andmatrix could be a source of detachment of HfC particles. It is difficultto confirm any signs for existence of abrasive wear as there were noparallel grooves on the track profile. Wear resistance dominantly relyon the dispersion of second phase and consolidation technique. See I.-Y.Kim, J.-H. Lee, G.-S. Lee, S.-H. Baik, Y.-J. Kim, Y.-Z. Lee, Frictionand wear characteristics of the carbon nanotube-aluminum composites withdifferent manufacturing conditions, Wear 267 (2009) 593-598,incorporated herein by reference in its entirety.

Wear Rate Study

FIG. 45 shows the comparison of preliminary results of specific wearrate for different loads for a sliding distance of 10m. These resultsshow a trend for load, sintering temperature and reinforcement contents.Wear rate increases with the increase of load and decrease with theincrease of sintering temperature and reinforcement contents. Sinceamount of wear was not significant due to smaller sliding distance, soanother set of experiments were performed to get a conclusive evidence.

FIG. 45 shows the wear behavior of monolithic sample sintered fordifferent temperature and composite containing 5 vol % HfC sintered at1800° C. A trend for increase in wear resistance can be observed fromthese bar graphs with the increase of sintering temperature and by theaddition of HfC contents. Wear resistance was found to decrease with theincrease of load. Since the wear rate order is low for 10 m slidingdistance, another set of experiment was performed for a sliding distanceof 100m.

FIG. 46 shows the comparison of wear rate of alloy and compositespecimens. It also provides the effect of sintering temperature on thewear resistance of alloyed specimen. Wear resistance is improved withthe increase of sintering temperature. This is in agreement with theFESEM analysis of the wear profile. This trend is also exhibited by mostmetallic materials. See J. Halling, Wear 35 (1975) 239; J. F. Archad, J.Appl. Phys. 24 (1953) 981; and A. J. Clegg, A. A. Das, Brit. Foundryman70 (1977) 333, each incorporated herein by reference in their entirety.An improvement in wear resistance was attained due to the enhancedconsolidation and improved hardness of alloy with the increase oftemperature as shown in FIG. 46. The wear resistance was furtherimproved for composite due to strengthening by the second phase hafniumcarbide. This fact is supported by the laws of wear i.e. the wear rateof a material is inversely proportional to the hardness of the softersliding component. See A. D. Sarkar, Wear of Metals, Pergamon Press,Oxford, 1976, p. 49, incorporated herein by reference in its entirety. Acomparison of relative density, microhardness and wear is presented inTable 4. The results showed that relative density, microhardness andconsequently wear resistance was increased for monolithic alloy with theincrease of sintering temperature. Wear resistance was found to improvewith the addition of 5vol % HfC in W-25 wt % Re matrix.

TABLE 4 Effect of sintering temperature and HfC on hardness and relativedensity Relative Spec. Temperature Microhardness density wear rateComposition (° C.) (Hv) (%) (mm³/N · m) W—25%Re 1500 360 92.2 2.64E−05W—25%Re 1700 396 96.5 4 77E−06 W—25%Re 1800 425 98.3 3.52E−06 5 vol %HfC 1800 495 95.2 1.00E−06

FIG. 47 shows the coefficient of friction (COF) of monolithic alloyvaries between 0.35 and 0.52 when sintered at 1500° C. with no evidenceof steady state for these conditions. FIG. 48 shows the coefficient offriction alloy specimen sintered at 1800° C. The COF varies between 0.63and 0.76 which can be due high contact forces between the surfaces.

Initially friction coefficient of the monolithic alloy increasedcontinuously and remained steady for the remaining interval of theexperiments.

Initial stage behavior can be attributed to a polishing process duringthe wear test, trying to establish a smooth wear track surface, byplowing away the surface asperities or roughness irregularities.However, COF remained steady at an average value of 0.39 for theremaining time interval. It is also evident that the composite'sfriction coefficient is not only lower but also has considerably lessfluctuation as compared to the monolithic alloy.

Example 7—Friction Stir Welding of High Melting Point Materials

The main objective of this section is to study the performance of anextruded W-25% Re pin tool for the FSW of ASTM A516 Grade 70 mild steelplates. The outcome of this part will provide base line parameters forthe feasibility of newly developed pin tool in FSW of steel in laterstages. The pin tool was designed by Edison Welding Institute (EWI) andmanufactured using extrusion by Rhenium Alloys Inc., in the USA. Mainemphasis will be put on the wear behavior of the tool. It will alsocover the effect of Friction Stir Welding (FSW) process parameters suchas tool rotation speed on the quality of the bead. Tool reactionsforces, microstructural features along with microhardness behavior wereinvestigated by varying tool rotational speed.

Experimental Set-Up

A set of single-pass partial penetration bead on plate (BOP) wereproduced using fully instrumented MTI (Manufacturing Technology Inc.)Model RM-1 friction stir welding machine. Proprietary W-25 wt. % Re toolwas used to produce the beads on ASTM A516 Grade 70 carbon steel platescontaining 0.25% C and 0.24% Si. ASTM A516 Grade 70 is widely used inproducing pressure vessels and heat exchangers.

A schematic diagram presenting the initial tool pin-shoulder dimensionsis shown in FIGS. 50A (Tool shape) and 50B (tool geometry in mm) alongwith FSW tool photograph. The tool consists of a simple tapered tool pinand a flat shoulder. This tool geometry was adopted for seal weldingapplications where full penetration is not required. In all tests,welding speed was ramped from 15 mm/min to a final velocity of 40 mm/min(15 mm/min (0-5 mm), 30 mm/min (5-15 mm), 40 mm/min (15-end mm), with acombination of tool rotational speeds of 800, 1500 and 2000 rpm. Speedwas gradually increased to avoid tool-pin fracture. The machine tiltangle was 0.1o and the dwell time was fixed to 3 seconds. Argonshielding gas was used during the friction stir welding of these beadsin order to avoid oxidation of the base metal, as well as to preventdeterioration of the tool pin surface.

Tool axial force and torque were recorded using the machine built-insensors as well as data recording system. Tool axial force and torquewere recorded using the machine built-in sensors as well as datarecording system. The data collected during the process was used to (1)study the effect of tool rotational speed on the dependent processparameters (i.e. tool axial load and torque), (2) find the relationshipbetween weld quality (surface integrity) and load profiles, and (3)estimate the heat input or line energy as shown in Eq. (5.1)

$\begin{matrix}{E = \frac{2{pNT}}{u\left( {1 + f} \right)}} & {{Equation}\mspace{14mu} 5.1}\end{matrix}$

Where f is the tool/workpiece heat ratio, N is the tool rotational speedin rpm, is averaged measured tool torque in N.m, and u is the tooltraverse speed (welding speed) in mm/sec. The tool/workpiece heatfraction f can be estimated using Bastier's model that describesone-dimensional steady-state heat transfer from a point heat sourcelocated at the interface of dissimilar metals presented in Eq. (5.2).See Bastier A, Maitournam Mh, Dang Van K, and Roger F., 2006, “Steadystate thermo-mechanical modeling of friction stir welding”, Science andTechnology of Welding & Joining, 11, pp. 278-288, incorporated herein byreference in its entirety.

$\begin{matrix}{f = {\frac{J_{T}}{J_{W}} = \sqrt{\frac{\left( {krC}_{p} \right)_{T}}{\left( {krC}_{p} \right)_{W}}}}} & {{Equation}\mspace{14mu} 5.2}\end{matrix}$

where k is the thermal conductivity, ρ is the density, and C_(p) is thespecific heat. Subscripts T and W indicates tool and workpiece,respectively.

In order to perform microstructural, microhardness and chemical analysisand investigate the effect of tool rotational speed on weld quality,fabricated beads were transversely sectioned in zones where finalwelding speed (40 mm/min) was achieved. Samples were mounted, grinded,polished and etched using 2% Nital solution to examine microstructuralfeatures and then perform microhardness measurements.

Optical magnifier was also utilized to study the nugget soundness of thedeveloped bead by revealing the processed zone and developed volumetricdefects if any. Optical microscopy analysis of the bead sections wasconducted at different magnifications in order to investigate themicrostructural details of the developed bead zones. Microhardness ofbead was measured using Vickers microhardness tester. The indenter loadwas set at 300 g for a period of 15 seconds. The hardness was measuredacross the bead and the depth from the bead's centerline. Moreover,spectroscopy analysis was performed to study the chemical composition inthe processed zone, and to understand the effect of tool rotationalspeed on tool diffusion wear.

The setup for workpiece fixture shown in FIG. 51 consists of backingplate for supporting the workpiece fixture, clamps and screws forholding the workpiece. Argon shielding gas was used to avoid theoxidation of the base metal at high temperature developed during theprocess.

The spindle and the tool were cooled by ethylene glycol coolant to avoidthe overheating of the tool assembly. Loads profiles were recorded atdifferent welding conditions. They were found to be strongly coupled tobead surface quality.

Example 8—Development and Characterization of FSW Bead-On-Plate (BoP)

Influence of Rotational Speed on Bead Quality and Reaction Loads

FIGS. 52A, 52B, 52C, 53A, 53B, 53C, 54A, 54B, 54C, 56, and 57 showdeveloped beads surface, with measured tool reaction loads at processconditions mentioned in the previous section. It can be seen fromrecorded tool loads that steadiness of tool axial force and torqueindicated a stabilized process, which is reflected on bead surfacequality (FIGS. 52A, 52B, 52C, 54A, 54B, and 54C). Tool axial force andtorque-displacement profiles were found to be strongly coupled to beadsurface quality. As mentioned previously, three different toolrotational speeds were investigated in this work. Three beads wereproduced at 800 rpm, shown in FIGS. 52A, 52B, 52C, 53A, 53B, 53C, 54A,54B, and 54C, while 2 more samples produced at 1500, and 2000 rpm shownin FIG. 56 and FIG. 57 respectively.

Although the same welding conditions were used for the first threebeads, some discrepancies in outcome findings were observed. They werebrought by material build-up on the tool shoulder and pin. In FIGS. 52A,52B, and 52C, the trend of the performed welding indicated a uniformpattern in the beginning. After the tool traveled a distance of 50 mm,fluctuations in torque and axial force accompanied with repeated patternsurface defect were noticed as indicated by the dashed circle.

After welding, the tool was examined by visually and Scanning ElectronMicroscope SEM as shown in FIGS. 55A and 55B. Visual inspection showslarge amount of steel deposition accumulated on the tool. The depositedsteel resulted in creating an eccentric tool profile. The tool withbuilt-up (or accumulated debris) was used to produce the second beadshown in FIG. 55C. Uniform (repeated) surface defects pattern wasobtained at every shoulder advancement (10 mm). These defects wereaccompanied by fluctuations in tool loads. Surface defect were similarto those formed at the end of the first bead in FIG. 52A, 52B, 52C.

Similar deposition was found on the tool after completion of the secondbead, which support that the formed surface defects were resulted fromthe tool built-up. To remove the deposited material (built-up), a smallbead was intentionally produced with larger over plunging depth, inorder to increase the process temperature. As a result, a clean toolsurface emerged. Using the clean tool a third bead was produced with asmooth and defect free surface manifesting a stable tool reaction loadsas shown in FIGS. 54A, 54B, and 54C.

The built-up of working material on the tool is extremely critical; asit modifies the tool profile that may result in reducing the weldquality. However, there are many parameters that control the workmaterial build-up; namely, the tool temperature and its surfacecondition. Enhancing the tool surface quality may reduce the problem ofthe build-up material.

Increasing the tool rotational speed from 800 to 1500 rpm and then to2000 rpm resulted in excessive fluctuations in tool reaction loads, withan increase in peak-to-peak axial force and torque values (FIG. 56 andFIG. 57). Fluctuations in torque could be attributed to the stick slipaction between the tool and the workpiece. Bead surface pattern producedat rotational speed of 1500 rpm was not uniform. It can be seen fromFIG. 56 that a smoother surface was generated at the end of the bead,with signs of excessive heat input (change in surface colors) at thebead sides (FIG. 56).

Moreover, the formed flash was larger and thicker as compared to otherperformed bead. At 2000 rpm, the generated bead surface shows repeatedpatterns of shoulder marks (FIG. 57). It may indicate that tool-machinedynamic characteristics play a role in identifying optimum processconditions and controlling joint and surface quality. From the obtainedtool load-displacement results and bead surface images, it can beconcluded that a more stable process could be achieved with bettersurface finish and smaller shoulder marks at lower tool rotationalspeeds. However, risk of material build up is more susceptible to occurin this case.

In welding, heat input or line energy is considered one of mostimportant parameters that may affect weld quality. It is alwayspreferable to perform welding at low heat input in order to avoidtemperature related problems and defect, for example, residual stresses.Using the average recorded tool torque and rotational speeds togetherwith Equations and (5.2), the line energy was calculated and presentedin Table 5 along with measured tool reaction loads. From Table 5, theaverage recorded spindle torque in all beads was found to fall in therange of 30 and 40 N.m. Varying the tool rotational speed did not resultin major changes in average torque, but it exhibited excessivefluctuations in tool loads, as mentioned previously. On the other hand,recorded axial force was found to decrease with increasing the toolrotational speed, where the average axial force was about 2500 N at 800rpm. However, at higher rotational speeds (1500, 2000 rpm) it wasdropped to 1000 N, but the heat input had increased.

TABLE 5 Average recorded tool reaction loads and calculated heat input.Tool Average Average Average Maximum Rotational Tool Axial Tool HeatTool Speed Force Torque Input Torque [rpm] [N] [N · m] [kJ/mm] [N · m]800 2443 41 1.8 41.7 1500 945 33 2.8 49 2000 1095 39 4.4 69

Influence of Tool Rotational Speed on the Microstructure

FIGS. 58A, 58B, and 58C show cross sections of the produced beadscaptured using optical magnifier. Beads were sectioned at locationswhere steady state condition was achieved (at welding speed of 40mm/min) in order to study the impact of the tool rotational speed aloneon the quality of the nugget. Optical magnifier and optical microscopewere used to examine the different developed microstructural zones,material flow, and volumetric defects. At 800 rpm a sound bead wasdeveloped with no visible defect as shown in FIG. 58A. Moreover,material flow contours were not visible. The welding process conditionsmay have resulted in a stable process for this bead. At higherrotational speeds, however, root defects in the nugget were observed, asshown in FIG. 58B and FIG. 58C. These defects may have been developedbecause of insufficient material flow at the tip of the tool pin. It canalso be noticed that flow lines are more visible, due to excessiveinterruption in metal flow, which was also reflected on measured tooltorque and axial force.

The nugget produced at low rotational speed has better smooth beadformation with small Heat Affected Zone (HAZ) due to stable andoptimized process conditions. On the other hand, the nugget developed athigh rotational speed (1500 and 2000 rpm) showed defects in the stirzones (SZ) below the tool pin toward the retreating side. The HAZ wasabout 0.9 mm in width for 800 rpm whereas it was found to be 1.8 mm for1500 rpm and 2.4 mm wide for 2000 rpm. Similar behavior was reported inanother study for welding speed used in tube-tubesheet welds where theyfound that increasing welding speed increases the size of the voidformed in the nugget.

In all beads, the HAZ in the advancing side was wider as compared to theretreating side. This could be attributed to the asymmetry intemperature distribution about the weld centerline, as the temperaturesat the advancing side are higher than those of the retreating side.Moreover, HAZ was also found in wider at high rotational speeds ascompared to lower ones. This was expected due to the increase in heatinput. It was also evident that there is change in color around the beadindicating larger HAZ area around the nugget.

Optical micrographs of the weld beads were examined at different regionsof the bead to assess the effect of the pin tool rotation speed onmicrostructural features. FIGS. 59A, 59B, and 59C shows the opticalimages of the base metal and friction stirred bead produced at 800 rpm.FIGS. 59A and 59B depict the optical microscopic images of the parentmetal, showing typical distinct phases of ferrite and pearlite. Theas-received base metal exhibited a microstructure of ferrite grainsapproximately 20 to 30 m in diameter ad smaller grains of fine pearlite.Grain size refining was observed in the SZ at all welding conditions andin all weld beads.

At 800 rpm tool rotation speed, equiaxed grains of almost 5 μm weregenerated as a result of thermomechanical actions during the weldingprocess as shown in FIG. 59C. When the rotational speed is increased to1500 rpm, the grain-coarsening took place, which could be attributed tothe increase in temperature and hence to increase in heat input. Furtherincrease in rotation speed to 2000 RPM resulted in higher temperaturesoftening and grain coarsening as compared to 1500 rpm conditions. Theobtained results also confirm that axial forces were found to decreasewith the increase of rotational speed resulting in a softening of theworkpiece which is attributed to grain coarsening. Intense plasticdeformation and heat input have resulted in recrystallization of thebead nugget at all investigated welding conditions. Therecrystallization of the SZ produced fine grain microstructure, which isof significant importance in post-weld hardness.

FIGS. 60A, 60B, 60C, and 60D show microstructurally distinct regionsnamely SZ along the bead centerline, HAZ around the SZ of the beads. NoTMAZ (Thermomechanical Affected Zone) was seen in the steel beads, whichis unlike to aluminium alloys. In FIG. 60A, ferrite and pearlite on theadvancing side were partially dissolved showing no distinct phases offerrite and pearlite due to higher temperature on the advancing side. Athigher rotational speed, however, more grain coarsening was observed inthe HAZ region adjacent to SZ as shown in FIGS. 60C and 60D. It isbelieved that temperature of the SZ is greater than that of the HAZ.Normally FSW processes of steel generates temperature around 1200° C.See G. E. Linnert, “Welding Metallurgy: Carbon and Alloy Steels”American Welding Society, 3rd ed., Vol. 2. pp. 312-325., 1967,incorporated herein by reference in its entirety.

The HAZ has no apparent deformation during FSW. Therefore,microstructural evolution in the various regions of the HAZ of the steelcan be compared to Fe-Fe3C phase diagram. It was also reported thatgrain coarsened region of HAZ experienced the highest temperature andresults of the some other researcher showed that temperature is above A3line temperature in the phase diagram meaning that the grain growth ofthe austenite will be inevitable. See T. H. Courtney, “MechanicalBehavior of Materials”. pp. 309-317. New York. N.Y.: McGraw-Hill, 1990,incorporated herein by reference in its entirety. Existence of any TMAZwas probably lost due to decomposition of austenite during cooling.

The microstural investigation of the SZ is not an easy task as it isdepicted in HAZ. It is also worth mentioning that microstural evolutionin FSW of steel is consistent with continuous cooling curve of the arcwelding of similar steels. See K. E. Hughes, K. D. Nair, and C. M.Sellars, “Temperature and flow stress during hot extrusion of steel”Metals Technology, Vol. 1(4), 161-169, 1974, incorporated herein byreference in its entirety. In fact, the HAZ bears only a thermal cycle,whereas the SZ experiences both thermal and mechanical cycles so bothprocesses will be considered to discuss the evolution of phases in theSZ. Actually during FSW of steel, dynamic recovery, dynamicrecrystallization, and metadynamic recrystallization in those regions,which are near the tool shoulder that bears large strain near thesurface. See W. M. Thomas, P. L. Threadgill, and E. D. Nicholas,“Feasibility of friction stir welding of steel” Science and Technologyof Welding and Joining 4(6): 365-372, 1999, incorporated herein byreference in its entirety. Large strain will increase the grain sizerefinement. Therefore, the SZ will experience an increase in grain sizerefinement as compared to other bead zones. For the remaining zonesincluding the bottom of the bead, they will experience lower strain andslow cooling rates resulting in an increase in grain coarseing.

Influence of Tool Rotational Speed on Bead Microhardness

FIG. 61 shows hardness of beads produced at different rotational speedsand constant welding speed of 40 mm/min. The hardness of the base metalwas 135±3 VHN. Microhardness of the SZ portion of the weld bead wasfound to be greater than those of the surrounding zones including the BMand the HAZ. This was expected due to the grain size refinement thatresulted from thermomechanical cycles in the SZ. The highestmicrohardness values were obtained at 800 rpm. Microhardness valuesranges from 275-310 VHN for 800 rpm. The increase in hardness valuescould be attributed to refined microstructure. Moreover, lowerrotational speed will generate lower temperature that will prevent graingrowth. At higher rotational speed (2000 rpm), however, grain coarseningoccurred and hardness decreased drastically. Hardness values recorded inthe SZ were found to vary from 200 to 240 VHN.

Similar observations were found by A. Pradeep and S. Muthukumaran. SeeA. Pradeep, S. Muthukumaran, “Two modes of metal transfer phenomenon infriction stir welding of low alloy steel plates” Proceedings of the 1stInternational Joint Symposium on Joining and Welding, Pages 305-312,2013, incorporated herein by reference in its entirety. Same trends wereobserved when hardness analysis was conducted along the depth of thebead as shown in FIG. 62. Microhardness values decrease from top tobottom of the bead. Higher hardness at top surface could be attributedto high strain and high cooling rates due to direct contact of toolshoulder with the workpiece. The remaining regions were having lowerhardness due to low strain and cooling rates.

Influence of Tool Rotational Speed and Travel Distance on Tool Wear

The W-25% Re tool under investigation has traveled almost 3.5 m, andperformed over 50 plunges. The tool pin plunging depth was initially1.60 mm, and after traveling for more than 3 meters, the plunging depthincreased to about 2.2 mm, this is due to changes in the diameters andlengths of the shoulder and pin. FIG. 63 shows the average values of theaxial force, traverse (welding) force, and torque on the traveldistance, i.e. the length of the beads made by the tool, when frictionsteel welding the steel plates at welding speed of 40 mm/min and spindlerotation of 800 rpm. The rotational speed is proportional to the toollife in service. The FIG. 63 indicates that traverse (welding) forceincreases with the tool life whereas the torque decreases with theincrease in tool life. This is due to the wear and mechanicaldeformation of the pin tool.

Table 6 shows chemical analysis of the base metal and weld bead used inthe present work. The analysis of the base metal shows that it hastypical mild steel composition. The chemical analysis was performedusing spectrometer. The analysis was conducted to study and understandthe effect of rotational speed on tool wear. The analysis of the basemetal shows that it has 0.0209 wt % tungsten.

FIG. 64 shows the graph for the corresponding Table 6. It is noticedthat increasing the rotational speed resulted in an increase in theTungsten contents of the weld bead. This phenomenon could be attributedto the rise in temperature with the increase of rotational speed. As aresult, more diffusion of tool material to the weld nugget is obtained.

TABLE 6 Spectroscopy analysis of the bead developed Specimens wt % C wt% Si wt % Mn wt % P wt % S wt % W Base Metal 0.244 0.241 1.09 0.0110.008 0.0209  800 RPM 0.261 0.200 0.912 0.0053 0.0031 0.0951 1000 RPM0.294 0.263 0.912 0.0077 0.0048 0.119 1500 RPM 0.250 0.225 0.917 0.00810.0057 0.4825 2000 RPM 0.264 0.238 0.909 0.0062 0.0048 0.855

FIGS. 65A and 65B show the initial tool tip pin and shoulder length.Overall, the tool had experienced severe conditions of fluctuatingforces and higher thermal input which resulted in wear out and plasticdeformation of the tool shoulder and pin. After covering a 3.5m distancein more than 50 plunges, the tool shoulder shank deformed significantlyas it was shortened by 0.8 mm as shown in FIG. 65A (right image) afterthe tool had traveled an accumulative welding distance of 3.5 meter. Itis well known that tool shoulder is mainly responsible for thefrictional heat generation. Therefore, it is subjected to severe thermaland mechanical harsh conditions of FSW. It can be seen that the lengthof the pin increased by 0.6 mm as shown in FIG. 65B (right image).

This change in length is due to the wear of the tool shoulder shank andthe effects of competing mechanical deformation mechanisms acting on thepin tool. By calculating the tool pin and shoulder shank volume for thetool in its initial use and after 3 m of welding, it was found that thetool pin shoulder volume dropped by 40%. Similar observations onexcessive wear of W-25% Re tool were reported by Shuaib et al. where itwas stated that abrasion wear was the main wear mechanism. In addition,creep may also take place due to excessive process temperature whilewelding steel.

Table 7 shows the base line parameters which may be consulted for thenewly developed nanocrystalline tool.

TABLE 7 Base line parameters obtained during the FSW of mild steelCategory Parameter/Process Category Parameter/Process Process FSW ToolMaterial W—25 wt % Re Spindle 800 RPM Materials ASTM A516 rotationalspeed Joined Grade 70 Spindle tilt 0.1° Type of weld Buttjoint anglejoint Dwell time 3 sec Control Displacement control Penetration 1.6 mmTorque 40 N · M depth Plunging rate 5 mm/min Forging force, Fz 2500 NWelding speed 15 mm/min to Traverse forces, 400 N 40 mm/min Fx and FyMicrohardness 286 ± 3 Thermal 45-50 Wm−1k−1 Hv conductivity

It was shown that, by using a small pin tool, defect free bead-on plate,with good surface finish and intensive grain refinement was achieved at800 rpm. Low rotational speeds are recommended for the newly developedtool. At high rotational speeds, grain growth occurs due to increase intemperature and consequently, axial force decreases. In HAZ, more graingrowth was observed at high rotational speed. Furthermore, thespectroscopy results showed the presence of W in the weld bead in therange of 0.0951 to 0.855 wt %. This is due to tool material transfer tothe weldment through competing wear mechanisms including diffusion,abrasion, and chipping. Shoulder of the tool was found to be more proneto wear compared to tool pin.

The wear of the tool shoulder could be attributed to abrasive wear(removal of material or thinning of shoulder height), adhesive wear (dueto stick slip action during the FSW process) and diffusional wear (asindicated by the spectroscopy analysis of the beads).

The wear resistant of the tool materials can be improved by synthesizingthem using a combination of novel techniques such MA and SPS. W-25% Recan be reinforced with HfC to improve the abrasive wear resistance. Toollife is depended on the tool degradation during friction stir welding.Metallurgical challenges which will affect tool life are crystallitesize of the matrix, homogenous dispersion of second phase in the matrix.Nanocrystallinity of the alloy is difficult to retain by conventionalconsolidating techniques as it involves longer sintering duration. Sothese challenges are difficult to address by using these techniques.Therefore, the author proposes using such techniques as mechanicalalloying and spark plasma sintering to prepare W—Re pin tools toovercome these challenges during welding steels.

Example 9—Friction Stir Welding of Thin Mild Sheet

Friction Stir Spot Welding FSSW tests were also conducted on thin mildsteel sheets by using disc shape nanocrystalline tool in order toinvestigate its feasibility for the process.

In order to evaluate the performance of the developed tool materialunder the harsh conditions of Friction Stir Spot Welding of steel,preliminary results are presented here. Fully alloyed W-25 wt % Re toolin cylindrical disc shape was used in Friction Spot Welding of 2 mm thinmild steel in order to confirm the soundness and suitability of tool forFSW of steel. AISI 4140 alloy steel holder was manufactured to grip thetool during the process. Fully instrumented MTI (ManufacturingTechnology Inc.) Model RM-1 friction stir welding machine was used toperform the FSSW tests. Welding speed was 40 mm/min, with toolrotational speeds of 400, 500 and 600 rpm for all the experiments. Themachine tilt angle was 0o and dwell time was fixed to 5 seconds. Argongas was used as shielding gas during the friction stir welding in orderto avoid the oxidation of the base metal as well as the surface of thetool. Temperature was recorded during the process using telemetry systeminstalled with the set up.

Initially, fully alloyed W-25% Re disc sintered at 1800° C. was used asa tool to join the thin sheets. The disc type tool and the tool holdershank were manufactured in the machine lab as shown in FIGS. 66A, 66B,66C, and 66D. A provision for the thermocouple in the tool holder wascreated near the tool surface for the measurement of temperaturedeveloped during the process. It was difficult to machine and grip thetool in machining set up as the tool was only 4 mm in height aspresented in FIG. 66B. Milling operation was successful on compositedisc containing 5vol % HfC however it offered a resistance to machiningdue the presence of hard HfC particle as shown in FIG. 66E. Compositesamples have also lower thermal conductivity as discussed in earliersection. Lower thermal conductivity of the tool can produce localizedheating of the workpiece resulting in hotspot in the welds.

FIGS. 67A, 67B, and 67C show W-25% Re tool along with the tool holderassembly. The test was performed using a rotational speed of 400 rpm andwith a welding speed of 40 mm/min.

During this process tool holder became soft due to frictional heatbetween tool and workpiece and finally it deformed in later stages asshown in FIG. 67D. It happened because the tool holder was not at allheat treated and does not have sufficient strength to sustain the loadduring the process. It is also evident that the process did not developthe sufficient heat input in the early part of the experiment and thedisc tool was slipped from the tool holder as shown in FIG. 67E,however, sufficient heat was generated at the end of the process toproduce a spot weld as shown in FIG. 67F. Maximum temperature wasrecorded around 750° C. during the process. However during this severdamage of the tool holder, tool showed excellent resistance to anydamage or wear except some slippage marks as shown in FIGS. 68A and 68B.Tool pin was 2 mm in height and it did not change after the experiment.

Friction Stir Spot Welding of Thin Mild Sheet

New tool holder was manufactured from AISI 4140 alloy steel. A heattreatment was performed on the tool holder to make it harder andstronger. Tool holder was heated at 845° C. (1550° F.) followed byquenching in oil. After hardening the alloy was given as temperingtreatment to get an appropriate hardness of 54 HRC. The tool wasretraced and machined with new geometry as shown in FIG. 69A. FrictionStir Spot Welding was performed on the workpiece with a rotational speedof 500 and 600 rpm as shown in FIGS. 69B and 69C. The process producedgood quality spot welds. Spot weld 1 was performed with a rotationalspeed of 600 rpm that shows a wider heat affected zone when compared tospot weld 2 produced at 500 rpm. There were no visible defects in weldzones. The tool showed an excellent resistance to any physical ordimensional changes after these tests.

Optical Microscopy

FIGS. 70A, 70B, and 70C show the optical image of the base metal andnugget junction of friction stir weld produced at 600 rpm. Parent basemild steel shows typical distinct phases of ferrite and pearlite. Grainsize refining was observed in the SZ which could be attributed to theincrease in temperature during the process as shown in FIGS. 70B and70C. The process temperature of SZ must be in austenitic zone of Fe-Fecarbide diagram in order to justify the grain size refining phenomenon.

FIGS. 71A and 71B show cross sections of the spot weld produced atrotational speed of 500 and 600 rpm respectively. Both conditionsproduced sound weld cross sections with no visible defect. It was alsoobserved that HAZ was in large spread for FSW spot welded at 600 rpmcompared to 500 rpm.

FIG. 72 shows the effect of rotational speed on the diffusion of Wcontents into the weld. It can be pointed out that when the rotationalspeed was increased, relatively increase in tungsten contents wasobserved in the base metal. This phenomena was attributed to theincrease in temperature with the increase of rotational speed and hencemore diffusion of the tool material to the base metal.

1-10. (canceled) 11: A method using ball milling to form aceramically-coated friction stir welding tool having a cylindrical,conical or triangular tip, and a shoulder, comprising: forming acomposite of a ball-milled tungsten-rhenium alloy and hafnium carbideparticles, by: ball-milling a tungsten-rhenium alloy for no more than 25hours to form a first powder, wherein a concentration of rhenium is inthe range of 20 to 30 wt %, relative to the total weight of thetungsten-rhenium alloy; mixing hafnium carbide particles with the firstpowder to form a second powder, wherein a concentration of hafniumcarbide particles in the second powder is in the range of 1 to 15vol %,relative to the total volume of the second powder; ball-milling thesecond powder for no more than 15 hours to form a third powder; andspark-plasma-sintering the third powder at a temperature of 1500 to2000° C. for no more than 10 minutes to form the composite of theball-milled tungsten-rhenium alloy and hafnium carbide particles,wherein the ball-milled hafnium carbide particles are dispersed withinthe ball-milled tungsten-rhenium alloy, and wherein the ball-milledtungsten-rhenium alloy has a crystallite size of no more than 100 nm;coating the composite with a lubricant; extruding the composite to formand extruded article; then at least partially removing the lubricantfrom the extruded article; then coating the extruded article with aceramic to form the ceramically-coated friction stir welding tool. 12.(canceled) 13: The method of claim 11, wherein a crystallite size of theball-milled tungsten-rhenium alloy in the third powder is in the rangeof 10 to 50 nm. 14: The method of claim 11, wherein the third powder iscompacted with a pressure of 40 to 60 MPa during thespark-plasma-sintering. 15: The method of claim 11, wherein thetungsten-rhenium alloy is ball-milled in an inert atmosphere with amilling speed of 200 to 300 rpm, wherein a ball-to-powder weight ratiois in the range of 6:1 to 10:1. 16: The method of claim 11, wherein thesecond powder is ball-milled in an inert atmosphere with a milling speedof 100 to 200 rpm, wherein a ball-to-powder weight ratio is in the rangeof 4:1 to 6:1. 17-20. (canceled) 21: The method of claim 11, wherein thelubricant is at least one selected from the group consisting of acolloidal graphite, a glass powder, a silica particle, a siliconadhesive and combinations thereof.